1

C h a p t e r 1

INTRODUCTION

CONTENTS

1.1 Preamble 2

1.2 Review of background theory 4 1.2.1 Mixtures of materials 4

1.2.1.1 Phase diagrams 4 1.2.1.2 Miscibility 5 1.2.1.3 Partial Miscibility and changes in miscibility 5 1.2.1.4 Phase separation, USCT and LCST 6 1.2.1.5 Solid Solutions 8 1.2.1.6 Eutectics and Monotectics 8 1.2.1.7 Thermodynamics of Liquid Mixtures 11 1.2.1.8 Kinetics of Phase Transformations 12

1.2.2 Hydrogen Bonding 13 1.2.3 Small organic molecules 14

1.2.3.1 Melting and crystallisation 14 1.2.4 Polymers 14

1.2.4.1 Thermoplastics and Thermosets 15 1.2.4.2 Basics including molecular weight and molecular shape 15 1.2.4.3 Amorphous polymers 17 1.2.4.4 Polymer crystallinity 18 1.2.4.5 Lamellar melting 25 1.2.4.6 Melting behaviour of semicrystalline polymers 26 1.2.4.7 Spherulites 26 1.2.4.8 Poorly and partially crystallising polymer types 27 1.2.4.9 Polymer-polymer miscibility 27 1.2.4.10 Polymer-diluent systems 28

1.2.5 Linear polyamides (Nylons) 29 1.2.5.1 History of polyamides 29 1.2.5.2 Strengths 30 1.2.5.3 Weaknesses 31 1.2.5.4 Chemical structure and polyamide types 31 1.2.5.5 Biological-polyamide parallels 33 1.2.5.6 Polyamide Hydrogen Bonding 33 1.2.5.7 Polyamide Crystallinity 34 1.2.5.8 Polyamide Crystalline Structures 35 1.2.5.9 Effect of polyamide Type and Segment Length on Crystal Form 36 1.2.5.10 Multiple crystalline forms are possible - Polymorphism 38 1.2.5.11 Effect of pressure on crystallinity, melting temperature and crystal form 38 1.2.5.12 Metastability 39 1.2.5.13 Brill Temperature 39

1.3 Relevant papers in the area to be covered in the research 40 1.3.1 Small molecule-small molecule 40 1.3.2 Polymers with small molecules 40 1.3.3 Blend interactions and hydrogen bonding 41 1.3.4 Polyamides and Polymers 42

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1.3.5 Polyamides and small molecules 43

1.4 The focus of the research project 45 1.4.1 Materials chosen 46

1.4.1.1 Polyamides 46 1.4.1.2 Small molecules 47

1.4.2 Sample blending and notation used for blends 49

1.5 Experimental Techniques Used 50 1.5.1 Thermogravimetric Analysis 50 1.5.2 Differential Scanning Calorimetry 51

1.5.2.1 Thermogram Overlays 53 1.5.2.2 Thermograms expected from thermal events 54 1.5.2.3 Assignment of “Spiky” Crystallisations to Carbazole or Phenothiazine 56 1.5.2.4 Phase diagrams derived from thermograms 57

1.5.3 Simultaneous Differential Thermal Analysis/Thermogravimetric Analysis 61 1.5.4 Fourier Transform Infrared Spectroscopy 62

1.5.4.1 General 62 1.5.4.2 Mid Range IR and hydrogen bond Interactions 64 1.5.4.3 Mid Range IR Frequencies of Interest 65 1.5.4.4 Mid Infrared Data Analysis for Blends 67 1.5.4.5 Near Infrared FTIR (NIR) 71

1.5.5 Small Angle X-ray Scattering 72 1.5.6 Solid state Nuclear Magnetic Resonance Spectroscopy 73

1.6 Structure of the Thesis 74

1.7 Summary 75

1.1 Preamble Linear polyamides, commonly known as Nylons, have a broad range of

commercial applications. They are used widely where their high melting

temperatures, high heat stability, toughness and abrasion resistance can be

used to advantage. Detailed knowledge of their material properties is needed

to optimise their processability and properties when blended with other

materials for a variety of purposes such, as the formation of membranes.

This thesis contributes to the understanding of high temperature solutions

of semicrystalline linear polyamides melt blended with two different

crystallisable small-molecule organic compounds, carbazole and

phenothiazine. It also covers the crystallisation processes that take place

during solidification to room temperature. It concludes that the major factor

affecting the resulting nano- and microstructure of the solid is the relative

crystallisation temperature of pure polyamide and compound.

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There has been much work on semicrystalline polymers blended with

amorphous polymers [1] . There is not a great deal in the literature on

semicrystalline polymers blended with semicrystalline polymers [2-6]. The

production of membranes with Thermally Induced Phase Separation (TIPS)

by using amorphous polymers with crystallisable small molecule diluents

has recently become an area of some interest for some people [7]. Little has

been reported on the area of semicrystalline polyamides melt blended with

small organic compounds [8]. Using the crystalline small organic molecules

affects crystallisation strongly because the small molecules are highly mobile

ahead of the crystallising polymer front.

The work therefore makes an important contribution to a somewhat

neglected area, particularly as it covers a range of polyamides with differing

repeat units, melting temperatures and crystal structures. An investigation

of these differences has led to a better understanding of high temperature

solutions of polyamides with some small organic molecules and of the

manner in which semicrystalline polyamides crystallise with normally highly

crystalline small molecules. It will also enhance our knowledge of complex

lamellar formation and small organic molecule crystallisation in a

semicrystalline combination of the two types of material.

The major tools for the investigation are Differential Scanning Calorimetry

(DSC), Thermogravimetric Analysis (TGA) and Fourier transform infrared

spectroscopy (FTIR) in Mid-range infrared and the Near-range infrared (NIR).

The original purpose of this research had been to gain a better

understanding of crystallinity in linear polyamides and an appreciation of

how hydrogen bonding affects polymer crystallinity. The two types of small

molecules used are potential hydrogen bond disruptors. The research has

led to the different focus for the work because the two diluents were shown

later by Fourier transform infrared techniques not to interact in the solid

state by hydrogen bonding with polyamides. It was, however, recognised

that scientifically interesting questions arose from some of the experiments

that had been undertaken. These had been found with material from the

first sample made in an ampoule for producing bulk blends from the melt.

These larger quantities were to be used for several characterisation

techniques requiring bigger samples than the few milligrams that could be

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produced in a DSC. The interesting results were the production of three very

separate sections of the sample with quite different colours, brick red, white

and fawn and of very different brittleness/hardness. Thermogravimetric

Analysis (TGA) showed that the weight percentage polyamide was different in

the samples and Differential Scanning (DSC) thermograms were also very

different for the three. These results are discussed fully in Chapter 3.

1.2 Review of background theory Much of the information covered in the next few sections may be found in

undergraduate textbooks on physical chemistry and materials science. It is,

however, still worthwhile to refresh our memories and briefly draw all the

basic concepts together to form the groundwork of the research environment

of the project. The topics are covered in a relatively superficial manner and

are meant only to lead the reader to the point where current research in the

field is discussed.

1.2.1 Mixtures of materials What happens when two different materials are put together in the same

environment at the same temperature and pressure will be explored. This

will provide a basis for what happens when polyamides are heated with

either carbazole or phenothiazine to the melt and then cooled down.

1.2.1.1 Phase diagrams

The main part of the research work covers mixtures of materials being

studied by DSC in solid and liquid states at pressures near one atmosphere.

In these conditions, the effects of ambient pressure are not intrusive in the

measurements. A simplification of the total equilibrium state is to only

consider solid and liquid phases, as will be done here. The maximum

number of degrees of freedom for a system with two components is three so

by fixing the ambient pressure at a nominal one atmosphere we can

effectively consider only the two variables, temperature and composition.

Equilibrium phase diagrams represent the different phases encountered in

latter parts of the text. The phase diagrams are plots of temperature against

composition with lines defining temperature/composition conditions that

lead to regions where there is a common phase. The liquidus is the line that

defines the lowest temperature where all material is in the liquid state. The

solidus is the line defining the highest temperatures where all the material is

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in the solid state. There are regions in addition to the all-liquid and all-solid

states where there is a liquid coexisting with one or another solid material.

The materials are considered to be in the ideal equilibrium conditions in the

discussion immediately below. The reality of the experimental conditions is

that the work has been done in non-equilibrium conditions and this will

cause some modification of the outcomes. The major experimental

technique used in the work was Differential Scanning Calorimeter (DSC) and

from the DSC output we can observe some of the physical changes that take

place with melting and crystallisation. The melting and crystallisation peaks

can be interpreted to give an understanding of the underlying phase

diagrams, although with the caveat that the observations may not lie at

exactly the phase boundaries for equilibrium phase diagrams.

1.2.1.2 Miscibility

It is instructive to first discuss liquids before discussing multiphase solids.

Two liquids are miscible in each other when the molecules of one material

are completely dispersed in another at an atomic level for all concentrations,

such as ethanol in water. It is common to talk of solvent and solute where

one material (the solvent) is in substantial excess. This becomes more

difficult in many of the cases to be discussed in this thesis because

concentrations ranging from just over 20% polyamide to 83% polyamide are

encountered as well as the pure materials. General usage of solvent is not

necessarily the best because there will be a number of cases where the

liquids are not in solution at specific temperatures. A more general word

that will be used is diluent, which is suitable for all cases here. In most

cases, the terms will not be raised and only the weight percentage polyamide

referred to.

1.2.1.3 Partial Miscibility and changes in miscibility

Partial miscibility occurs when one liquid can only be added to another to a

certain limit and then will not dissolve further. The two materials will

separate out from one another into two layers if there is a difference in

density or into droplets/blobs of one material in the other if there is little

density difference. It should be noted here that there will be at least a small

amount of material A dissolved in material B and vice versa, even if the

materials are essentially immiscible, such as water and oil.

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Liquid combinations that are immiscible at one temperature can often

become miscible at other temperatures. For example if a mixture of phenol

and water is heated to over 65 0C it becomes miscible.

1.2.1.4 Phase separation, USCT and LCST

0 50 100 Percentage of Material B

Tem

pera

ture

Upper Critical Solution Temperature

(UCST)

Solution of A and B

A and B not in stable solution

Figure 1-1 Example of phase diagram with Upper Critical Solution

Temperature.

0 50 100 Percentage of Material B

Tem

pera

ture

Solution of B in A

Solution of A in B

A and B not in stable solution

A and B not in stable solution

Figure 1-2 Example of Phase diagram with immiscible region but no UCST or

LCST.

0 50 100Percentage of Material B

Tem

pera

ture

Solution of A and B

Lower CriticalSolution Temperature

(LCST)

A and Bnot instable

solution

Figure 1-3 Example of phase diagram

with Lower Critical Solution Temperature.

0 50 100 Percentage of Material B

Tem

pera

ture

Solution of A and B

A and B unstable and

spinodally decompose

Binodal line

Spinodal line

Metastable regions

Figure 1-4 Example of binodal and spinodal lines in a phase diagram.

The phenol/water case is an example of an Upper Critical Solution

Temperature (UCST) where there is a maximum temperature at which the

materials are not completely soluble. A typical phase diagram is given in

Figure 1-1 The UCST does not have to lie at the centre of the concentration

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range but is often very strongly towards one or the other side of the phase

diagram. There are also cases with some materials where there is a Lower

Critical Solution Temperature (LCST) and once the temperature has been

raised sufficiently the two materials begin to separate into separate phases.

That can be seen in Figure 1-3. Another case is shown in Figure 1-2, where

there is no upper or lower critical temperature but a region in mid

concentrations where the materials are insoluble. Utracki [9] in his work on

polymer/polymer miscibility states that UCST is more common in general

with solvent-polymer and LCST with polymer-polymer systems. There are

regions in composition-temperature space where the two materials cannot

exist stably in a single, miscible, phase. This line is called the binodal curve.

A metastable condition is often reached where the two materials still coexist

without separating if the density of the two materials does not differ

markedly. Another region exists within the binodal curve where the single

phase nature of the liquid becomes completely impossible. That inner

boundary is called the spinodal. Inside it the two materials will begin to

phase separate spontaneously. An example showing the spinodal in a phase

diagram is given in

Figure 1-4. Spinodal decomposition takes place throughout the mixed liquid

with very small volumes segregating themselves into like kinds of materials.

This is energetically unfavourable because of the high interfacial surface

area. Over a period of time, volumes of like material touch each other and

reduce surface area by coalescing. The entities of each material

progressively become bigger, as in Figure 1-5.

Figure 1-5 Ripening over time of small spinodally decomposed regions on the left to larger ones on the right.

This can happen by Ostwald ripening where domains at a greater radius

than some critical value grow at a faster rate by diffusion from the

surrounding medium. It can also happen by coalescence of droplets or by

hydrodynamic effects [10].

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The example shown has near equal amounts of each material but, when the

two materials are there in different proportions, droplets of one material can

exist in a matrix of the other material. Concentration changes encountered

by addition of one material or crystallisation can lead to a phase inversion

where the dominant body can become the droplets in the other.

The two materials are in a metastable situation if they are quenched to a

position on the phase diagram between the binodal and the spinodal.

Statistical density fluctuations often lead to phase separation by nucleation

and growth when the mixture is in the metastable region.

1.2.1.5 Solid Solutions

Solids can also form solutions in the same way that liquids do. It is the

ability of the solids to mix in all proportions of the basic materials that is the

criteria for a solid solution. In this case, though, the phase of the solution is

a solid rather than a liquid. An example of this is copper with gold.

1.2.1.6 Eutectics and Monotectics

There are many cases where solids are not significantly soluble in each

other but the liquids become soluble when the temperature is raised

sufficiently. One example of this is common solder used in electronics

where a eutectic is formed. Eutectic is Greek for “easy melting”. A

eutectic reaction is defined [11] to be:

“An isothermal, reversible reaction between two (or more) solid

phases during the heating of a system, as a result of which a single

liquid phase is produced.”

In these cases the phase diagram is similar to Figure 1-6. We see six regions

in the figure. The first has both together as a solution (in the melt). The

second and third regions α and β are solid solutions with one or the other of

the materials in virtually pure solid form with a small amount of the other

material dissolved in it up to the solubility limit. The fourth and fifth have

one of α or β as excess solid in equilibrium with the solution. The sixth is

where materials α and β are together in solid form. This last usually has a

finely divided matrix of α in β (or vice versa) with an overall concentration of

the eutectic composition. Within that are larger domains of any excess α or

β form.

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Consider solidified material after cooling from a molten mixture of materials

A and B and having a concentration and temperature defined by point c in

Figure 1-6. Material A is in excess so the solid will comprise nearly pure

material A inclusions (of phase α) with the same composition as m solidified

within a matrix of a solid eutectic mix of A and B (m and n with overall

composition of e). The solid will reach d as it is heated. At that point, the

eutectic portion will melt at the eutectic temperature (Te) into a liquid of the

eutectic composition, leaving the inclusions of α in equilibrium with the

liquid.

100

120

140

160

180

200

0 50 100

Percentage of Material B

Tem

pera

ture

(

0 C)

solid α + β

liquid

n

α + liquid

e

β + liquid

α β

h

f

m

d

g

j

k

c

Figure 1-6 A simple phase diagram showing eutectic formation.

A further increase in temperature will see some of the solid inclusions

changing composition along the line f-g as A dissolves into the liquid. This

causes the composition of the liquid to move along the line e-h as the

material α dissolves into it. The strong move of the liquid to the left with

increasing temperature means that a considerable amount of α is dissolving

into the liquid. Eventually the composition of the liquid will become the same

as the original proportions of the two materials in the solid at the time the

last of the α phase of composition g dissolves into the liquid. Further heating

maintains the composition at c-h and the liquid moves on the phase diagram

in the direction of j.

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Consider the alternative of a solution having a concentration and

temperature defined by point j in Figure 1-6. The state of the solution will

reach h as it is cooled. At that point, material A will begin to crystallise out

in nearly pure form as phase α with a composition given by point g across

the tie line linking compositions in equilibrium at that temperature. The

removal of phase α by crystallisation will naturally increase the relative

concentration of B in the solution as the temperature is lowered slightly.

The solution will thus follow the curved line towards point e as the

temperature is lowered further with continuing crystallisation of phase α.

The material crystallising will vary slightly in composition following the line

g-f. The lowest temperature where the liquid can coexist with solid is at

point e. The liquid cannot exist below the eutectic temperature so the

remaining liquid (of the eutectic composition) will crystallise at that point in

the cooling process. The final solid will incorporate solid, nearly pure A in a

matrix. That matrix is phase separated A-rich and B-rich micro domains

overall having the eutectic composition. On average, the solid will obviously

have the composition proportions of the two materials in the original

solution.

The above descriptions for the phase diagram are for equilibrium at all

times. Heating under practical conditions may mean that the composition of

the α inclusions may not have time to change from f to that of g. There may

be kinetic delays meaning that at faster heating rates some steps take place

a little later (at higher temperatures). The situation is more complicated

where we start with two powders A and B placed together. The powders will

reach the eutectic temperature where the points of contact between the two

types of powder will start to dissolve those grains of the powders. This will

continue until all of B powder is consumed, leaving pure powder A (in this

case) in the liquid of the eutectic composition. Further increases above the

eutectic temperature will result in the dissolution of powder A into the liquid,

moving the composition of the liquid along line e-h, as previously. The

practical implementation of eutectic formation from powders may result in

further delays than when starting with an α-in-eutectic solid. We will see

later that polymers often partly crystallise forming nanometre-thick

crystallites that tend to exclude other molecules. It may be expected that the

melting of a previously solidified mixture incorporating a semicrystalline

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polymer will behave in a fashion intermediate between two powders melting

and that for eutectic mixes of small molecules or metals.

The phase diagram seen in Figure 1-6 is one with a simple eutectic

relationship between the materials. Many, more complicated, types of phase

diagrams are found in practice with various material combinations. A

simpler phase diagram occurs with the side regions disappearing when the

solubility of one solid material in the other is totally insignificant.

A monotectic reaction is similar to a eutectic reaction but here a solid and a

liquid solidify (reversibly) from monotectic liquid. The compositions of the

solid and liquid are both different from that of the originating liquid. It is

possible to have multiple very small regions of liquid dispersed within a solid

matrix as a result of a monotectic reaction. Those liquid domains can then

solidify at lower temperatures.

IUPAC [11] define a monotectic reaction as:

“The reversible transition, on cooling, of a liquid to a mixture of a second liquid and a solid.”

1.2.1.7 Thermodynamics of Liquid Mixtures

Liquid mixtures, like other systems, are characterised by normal

thermodynamic parameters such as Gibbs free energy G, internal energy U,

enthalpy H, entropy S and volume V. The values found for real mixtures are

not the sums of the values of the pure constituents. For example, mixing a

volume of one liquid with an equal amount of another liquid will not give

exactly twice the volume of the first material. The same applies to the Gibbs

free energy. The difference between the actual free energy G and the sum of

the Gibbs free energies of the pure components Gi is the free energy of

mixing ∆Gmix. Similar comments apply to entropy and enthalpy with the

convention that the value for the mixture takes on the sign of the

subtraction of the sum of the components from the value for the real system.

This means. ∆Ymix = Y - (Y1 + Y2 +…+Yn), where Y is a thermodynamic

parameter and the values of Y with subscripts are those for the n pure

materials in the system.

There is a partial molar property for any of the above in a system defined as

the partial derivative of that property with respect to the number of moles of

one constituent when temperature, pressure and the number of moles of all

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the other components are kept constant. The partial molar Gibbs free energy

is the same as a parameter called the chemical potential of that constituent.

It is usually given the symbol µi.

These chemical potentials of the constituents are the quantities that

determine phase equilibria and from them we can derive a range of other

parameters. It is worth mentioning here that the chemical potential of a

pure material will be the same as the molar Gibbs free energy of the material

at the particular temperature and pressure of interest, ie. µi0 = Gi0.

This discussion is general to mixtures of liquids but will play a part in the

discussion later of the Flory-Huggins theory as it relates to polymers in

solutions.

1.2.1.8 Kinetics of Phase Transformations

Most transformations from one state into another do not take place

instantaneously because of impediments to the changes. Often energy

barriers related to the phase boundaries have to be overcome for the

molecules to be able to re-arrange themselves. There is usually a nucleation

stage followed by a growth stage. The whole process is time dependent.

The nucleation is the formation of stable microscopic particles of the new

phase in the originating phase.

This is followed by the growth of new material onto the nuclei. The growth of

the new material proceeds by diffusion into the old phase. It occurs until all

the volumes of new phase impinge on each other making the system wholly

the new phase.

The time taken for the change to take place is termed kinetics and is

obviously important for production processes. The rate at which the volume

of material changes from one state into another will be dependent upon how

much of the old state remains if we hold the temperature constant. The

outcome is an “S” shaped curve Figure 1-7 below that is described by the

Avrami equation in Equation 1-1 [12-14].

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0

20

40

60

80

100

0 40 80 120 160 200

Time (sec)

Per

cent

age

Pha

se

Tran

sitio

n

Figure 1-7 A typical Avrami plot for extent of crystallisation taking

place.

Y = 1-exp(-ktn) Eqn (1-1)

where Y is the volume fraction of crystalline material formed by time t at

constant temperature, and k is a variable dependent on temperature. The

exponent n should be an integer between 1 and 4, depending on the model

used, according to the original theory. Nowadays it is regarded as a variable

to match the data.

The Avrami approach is a simple one that has found application across a

wide variety of phase transitions. It is often used for crystallisation of metals

but is used for polymers [15, 16] as well.

1.2.2 Hydrogen Bonding We will now look at the hydrogen bonding that plays a strong part in the

behaviour of polyamides and the different types of bonding encountered in

chemistry.

Most people are familiar with ionic and covalent bonds. Ionic bonds take

place with the complete transfer of electrons from one atom to another and

are very strong. Covalent bonds are directed between two atoms such as

C-C or C-N and are also strong. Covalent bonds have energies in the order

of 300kJ/mol. The much weaker van der Waals forces are of the order of

1kJ/mol and are non-specific in direction.

Hydrogen bonds are intermediate in strength (around 30kJ/mol) and act

between hydrogen atoms and two other electronegative atoms, usually from

the group oxygen, nitrogen and the halides, particularly fluorine. They are

essentially electrostatic in nature. They act when hydrogen has been

covalently bonded to one of the above highly electronegative group that has

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drawn some of the charge from the hydrogen atom. This makes the

hydrogen atom partially positive in charge. An atom in another molecule or

another part of the same molecule that is also electronegative will be weakly

attracted to the hydrogen atom, forming a hydrogen bond (or hydrogen

bridge).

Probably the most important case of hydrogen bond formation is with water

where an O-H from one water molecule is bonded to the O of another water

molecule. Hydrogen bonds are continually forming and reforming, even in

water near 100 0C. The reason the water in our own bodies does not

evaporate at sub-zero temperatures is the hydrogen bonding that provides

an energy barrier to evaporation. It is also hydrogen bonding forces that link

the peptide groups of DNA into a double helix. We will see later that

hydrogen bond formation is implicit in the physical properties of the

polyamides (Nylons) of this research.

1.2.3 Small organic molecules 1.2.3.1 Melting and crystallisation

Small organic molecules in the solid state are arranged in a very regular,

symmetrical manner, held in place by van der Waals and possibly hydrogen

bonding forces. These forces will be stronger if the molecules can fit closely

together with as many atoms of one molecule as close as possible to atoms of

the next. The covalent forces holding each molecule together are much

higher than the inter-molecular forces. The atoms gain in vibrational and

rotational energy as the temperature is raised until the structure breaks

down suddenly and a disordered liquid state results. The reverse occurs as

the temperature is lowered and the molecules can nestle together in an

ordered structure. A better physical fit between the molecules will result in

stronger van der Waals forces between the molecules and a higher melting

temperature.

1.2.4 Polymers The Greek word poly means many and the word meros means part.

Polymers are macromolecules (very large molecules) made from the

combination of a large number of smaller repeating molecular units

(monomers) to create a larger molecule. They occur naturally as in DNA or

silk and can be manufactured synthetically as in Nylons and cured epoxy.

This research covers synthetic homopolymers made from only one sort of

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repeat unit unlike copolymers where there are polymer sections of differing

types within the one molecular chain. They can be structured as long

chains, with or without sidechains, networks, or dendrimers. This work is

on long linear polymers without sidechains.

1.2.4.1 Thermoplastics and Thermosets

Polymers are of two types, thermoplastic and thermosetting. Thermosetting

polymers result from the in-situ reaction of smaller molecules where covalent

crosslinkages form between the molecules during polymerisation, resulting

in a network. Subsequent reheating of the solid will not allow the material

to liquefy once the polymerisation reaction has been performed. Eventually

degradation occurs with the application of further heat. Thermoplastic

polymers can undergo multiple cycles of the polymer softening and becoming

liquid with heat and hardening on cooling. The polymer chains gain

sufficient vibrational energy to break the weak van der Waals forces between

the polymer molecules during this reversible process although this usually

takes longer than with small molecules because of the greater number of

molecular interactions involved with these large molecules. This type of

polymer can be processed in the melt by injection moulding, casting, blow

moulding and spinning to form solids of the required form on cooling. The

research here is on thermoplastic polymers.

1.2.4.2 Basics including molecular weight and molecular shape

The number of repeat units in a polymer molecule affects the size of the total

molecule in solution or the melt. This, in turn, affects the viscosity in

solution and the melting temperature. Commercial polymers often have

molecular weights of 20 to 50 kilodaltons.

Reactions to make polymers from monomeric units normally do not mean

that all molecules form at exactly the same molecular weight. There is

usually a distribution of molecular weights from a polymerisation process.

This means that the physical properties of a polymer such as viscosity are

an average over all chain lengths represented in the sample. The degree of

polymerisation, n, will be a distribution with a number average, Mn, that is

less than the weight average, Mw. The ratio of the two is called the

polydispersity and is a measure of the broadness of the molecular weight

distribution.

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Polymer chains have the opportunity to become entangled, spaghetti-like,

when in solution or the melt if they are sufficiently long. Viscosity is

increased markedly once entanglement has set in and the dynamics of

crystallisation are also altered [17, 18]. The number of repeat units before

entanglement becomes an issue is different between polymers and depends

on whether the polymer is a plain linear chain or has side chains, and on

other characteristics of the repeat units.

The distance between atoms of different polymer chains is a balance between

attractive van der Waals forces and Born repulsion between the clouds of

electrons surrounding each atom. The bond lengths between covalently

bonded atoms in the one molecular chain attempt to remain at their

equilibrium distances. At the same time, the bonds try to stay at their

optimum angles. Every atom in a polymer chain attempts to find an

energetically favourable position for itself under the constraints of bond

angles and interatomic distances. The application of more heat to a system

results in greater vibration of atoms around their optimum positions. We

will see later that the multiple forces acting on an atom can be utilised in

Fourier transform infrared techniques to characterise the environments of

atoms by their frequencies of vibration.

Polymer chain molecules are not straight. Usually the bonds are at preferred

angles other than 1800 and, unless sterically hindered by some of the atoms,

are able to rotate when in the melt or in solution. This results in a three-

dimensional “random walk” if the path in space is followed from one atom to

the next as displayed in two dimensions in Figure 1-8. The length of the

chain from one end to the other can be seen to be much greater than the

end-to-end distance along the straight line A-B. The size of the molecule can

be characterised statistically for a given situation with the radius of gyration

as a measure of how large the molecule is. That is determined by the mass

average of the square root of the squares of distances of the atoms from the

centre of mass of the molecule. A flexible molecule that is in thermal motion

requires that a time average be taken over all configurations.

The thermal vibrations in a melt at high temperature will tend to result in a

larger radius of gyration. The size of a polymer chain in a poor solvent will

be much smaller than with a good solvent because the chain segments tend

to keep to similar chemical environments. They retract to a smaller volume

17

to exclude unfavourable solvent molecule interactions. This also occurs with

proteins in an aqueous environment where the hydrophobic portions bury

themselves at the centre away from the solvent and the hydrophilic portions

extend into their watery surrounds.

Figure 1-8 Random walk between A and B, the ends of the polymer chain.

A linear polymer chain, such as with the Nylons of this text, will have a

larger radius of gyration than one of the same molecular weight but with

branches emanating from the main chain.

1.2.4.3 Amorphous polymers

Polymers can solidify into the amorphous state where, firstly, the

longitudinal motion is locked in, making the polymer solid but rubbery.

When the material reaches the glass transition temperature (Tg), it becomes

like a glassy, supercooled liquid as it is cooled further. Here, the motion

allowed by molecular vibrations and rotations becomes much more

restricted. Thermal conductivity, dielectric constants and mechanical

properties undergo significant changes at similar (but usually not quite

identical) temperatures. The glass transition temperature (Tg) for a particular

parameter is the point at which half the step change has taken place.

Polymers are very stiff below their Tg but they become rubbery as the

temperature is raised again. Eventually they become viscous liquids that

quickly thin further as the temperature is progressively raised. The process

is completely reversible unless the temperature has been raised so high that

the polymer has degraded.

An example of the viscosity behaviour as temperature is raised for an

amorphous polymer is shown in Figure 1-9.

18

Temperature

Visc

osity

|

Tg

Figure 1-9 Typical effect of temperature on viscosity for amorphous thermoplastic polymers. Viscosity is extremely high below the Glass Transition temperature Tg and rapidly drops with increasing temperature.

The hole theory of liquids requires there to be minute voids between the

molecules that allow them to move from one position to another. If we

extend this to polymers we must recognise that the molecules of a polymer

chain have to move cooperatively. This requires a minimum void size to

allow chain segments to move from one location to another. The free volume

will increase rapidly with temperature above this critical temperature. The

free volume of a polymer will remain relatively constant below this

temperature as molecular motion is frozen.

1.2.4.4 Polymer crystallinity

The discussion in this and the following four sections is only meant to

provide a general background for discussion of various aspects of polymer

crystallinity in the main chapters dealing with specific polyamide-diluent

combinations.

Over fifty years ago it was found that approximately three quarters of

different types of polymers are able to also enter a partially crystalline state.

This can happen for those types of polymers if the cooling conditions are

slower, or if solid amorphous polymer is taken through an appropriate

thermal history that allows a solid-state crystallisation to take place. The

degree of crystallinity will depend on the thermal and mechanical history of

the sample and can range from zero to 90%. It is this crystallisation that

adds to the mechanical stability of many manufactured plastic articles.

Generally, as the cooling rate during crystallisation increases, the percentage

in the amorphous state increases and crystallinity decreases. Polyethylene

19

and polyethylene oxide tend to have very high crystallinity and others have

much less, ranging down to those that cannot be crystallised.

Polyamide-4,6 being studied here can crystallise up to 70% by volume under

favourable conditions.

Lamellae are crystalline regions within the overall amorphous polymer. The

order caused by polymer chains (and parts of polymer chains) aligning

themselves means regularity in the structure of atoms, allowing Bragg

reflections to be seen with X-Rays in the manner seen with crystalline

mineralisations.

Whether a crystallisable polymer solidifies purely with an amorphous

structure or with a certain extent of crystallisation will depend upon a range

of parameters that can also affect the crystallographic form. The thermal

history as well as the molecular weight will also play a strong part. A final

crystalline state, probably metastable, will depend on nucleation states and

entropic barriers. Other factors that can affect the way in which the final

crystal form develops are the degree of undercooling, recrystallisation and

the lamellar thickening or thinning mentioned below.

The crystallisation takes the form of lamellae (platelet like structures or

crystallites several micrometres across and 5 to 10 nm thick) as seen in

Figure 1-10.

Figure 1-10 Keller’s diagram [19] for the laying down of folded polymer chains along the edge faces of lamellae

The larger, flat surfaces are called the basal planes and the thin surfaces

along the edges (and joining the basal planes) are called edge faces. Long

polymer chains are generally considered [19] to fold backwards and forwards

into place across the edge faces of lamellae crystallising from the melt or

solution. These energetically unfavourable assemblies come about mainly

because of the attempts of long polymer molecules to rearrange themselves

20

into energetically more favourable structures with greater order. They are, of

course, restricted in their ability to “reptate” like snakes into ideal positions

in reasonable timescales. Kinetics plays an important role in the perfection

attained.

Shorter polymer chains can crystallise in an extended chain form where

molecules line up together, side by side, without chain folding. This is a

lower energy configuration because there are no folds necessary. For

example, PEG usually forms folded chains in the lamellae only when the

molecular weight exceeds approximately 4,000 Daltons [20, 21]. The

extended chains become difficult to lay side by side when they are too long.

Layers are added on the edge faces to build up the thin lamellae from their

centre with chain folds in a consistent manner dependent upon steric effects

with the atoms and energy minimisation. The folding is driven by kinetic

factors because the initial nucleus has polymer chains locked into the folded

form.

The lamellae formed in solution are usually more perfect than those formed

in the melt because there is more opportunity for polymer chains to easily

orient themselves correctly by displacing the smaller solvent molecules.

Later growth of the lamellae by secondary nucleation of chains onto the edge

faces of the crystals continues the original folding form but often the

thickness of the lamellae vary as they grow bigger. It is common for

polymer lamellae to thicken if they are later annealed for some time at

temperatures somewhat below the melting temperature. That occurs by

polymer chains reptating like snakes in the lamellae to produce the more

thermodynamically stable thicker configuration. The ends of molecules

withdraw from their initial place in the lamella during the reptation process.

It has been found [22] that lamellar thinning also occurred with some semi-

rigid polymers, including polyamides.

The lamellae can often form in different crystallographic structures.

Generally the lamellar thickness increases with molecular weight and the

melting temperature also increases.

The lamellar thickness generally depends on the temperature of

crystallisation in polymers. The thickness of a lamella is reduced when

crystallisation takes place at a larger undercooling. The lamellae do not

have time to form in thicker, more energetically favourable forms when

21

driven by high supercooling. The significant surface energy tied up in thin

lamellae makes them less stable, in general, than thicker ones. This leads to

the melting temperature of thinner lamellae being reduced below that of an

infinitely large crystal. The melting temperature approaches an asymptotic

value as the molecular weight tends towards infinity.

Gibbs and Duhem were able to relate the equilibrium melting temperature,

Tm0 to the measured melting temperature Tm using the lamellar thickness l,

the enthalpy of fusion ∆Hf and σe from the slope of the graph of Tm against

1/l. Hoffman and Weeks [23, 24] took this further by eliminating the need to

know the lamellar thickness with the use of plots of Tm against the

crystallisation temperature Tx under the same ramp rates. Measured values

for the pair are extrapolated to the line Tm = Tx and that gives the equilibrium

value 0mT .

The Hoffman-Weeks approach has been extended with linear and non-linear

extrapolations by Marand, Xu and Srinivas [25]. On the other hand Welch

and Muthukumar [26] believe that a reliable estimate of equilibrium melting

temperatures cannot be obtained by this method.

We must consider two aspects for crystallisation to take place, the

nucleation of lamellae formation and the kinetics of lamellar growth. There

is the primary nucleation of lamellae as a first stage and then and then the

growth stage with secondary nucleation.

The primary nucleation can be likened to atoms in a gas condensing with

lowered temperature to form small groups with a high surface area to

volume ratio. This situation is unfavourable with much energy tied up in the

surface compared to the free energy gain by condensation to a liquid. The

group will dissociate unless sufficient atoms can simultaneously coalesce to

make a nucleus that can grow. That is because the increase in volume

produces a lower energy than the increase in surface area. The process

relies on statistical fluctuations at high temperature. These intermediate

groups then produce nuclei that can grow further. The formation of lamellar

nuclei in molten polymer or in solution, are generally regarded as analogous

to the gas-liquid condensation described above. An embryo lamella accretes

and loses adjoining sections of polymer chain in a dynamic manner until it is

large enough that the free energy gains outweigh the surface energy effects.

22

It is able to become a lamellar nucleus for further growth. This is a type of

situation with a hurdle to overcome where the Avrami approach is

applicable.

The major theory or model put forward to cover the secondary nucleation and

growth of lamellae once they have nucleated and begun to grow is that of

Hoffman and Lauritzen [27, 28]. In that theory there is assumed to be a flat

existing substrate. A new polymer chain lays down on the flat surface of the

edge face beside the previous chain and becomes attached to both. The

question then arises as to when it folds. An attempt to maximise the contact

area and bonding for the new chain will lead to longer lengths between folds

and result in thicker lamellae. The necessity to achieve this quickly when

there are strong driving forces towards crystallisation means a shorter length

between folds is desirable. The final lamellar thickness is thus dependent

upon the interplay between these opposing tendencies. The outcome is that

larger undercoolings result in thinner lamellae.

Often crystallisation also takes place more slowly behind the main

crystallisation front on the lamellae. This is called secondary crystallisation

and results in increased densification of the solid, particularly because the

slower rate of crystallisation results in more perfect (and denser) crystalline

regions. Diluent in the melt blend systems studied here will mean that

secondary crystallisation effects for the polymers will be promoted due to

dilution but be reduced by lower viscosity of the uncrystallised

polymer-diluent material. The outcomes cannot be predicted and would

require other techniques such as time resolved SAXS measurements to

determine. An analogous situation can be seen to occur for the diluent with

mini spikes occasionally being seen well after the main diluent

crystallisation peak such as with 65PA46Car in Figure 3-14 where there is a

small diluent crystallisation spike 30 0C lower than the main carbazole

crystallisation peak. Androsch and Wunderlich [29] showed with studies on

poly(ethylene-co-octene) using Temperature Modulated Differential Scanning

Calorimetry (TMDSC) that secondary crystallisation occurred with a delay of

5 min after primary crystallisation when cooling at 10 0C/min.

There is, however, much controversy over the steps that take place in going

through from the supercooled melt to the formation of crystals.

23

Olmsted et al. [30] suggest that there is liquid-liquid spinodal decomposition

taking place prior to the formation of crystals. The spinodal decomposition

was detected by Small Angle X-ray Scattering (SAXS).The actual crystal

formation was detected by Wide Angle X-ray Diffraction (WAXD). They base

this in part on work by others eg Ezquerra [31] on a variety of polymers

where SAXS signals are seen to increase and partially decay before the

WAXD signal appears in simultaneous SAXS-WAXD. They propose that

there are statistical fluctuations in density and entropy (linked) which result

in spinodal decomposition of the molten polymer into denser and less dense

phases. It is the more ordered, dense regions (detected by SAXS) that later

crystallise into the lamellae detected by WAXD. The direct experimental

results are supported by Monte Carlo simulations by Toma, Toma and

Subirana [32] where they investigate the formation of a compact globule

state with a lamellar conformation prior to the creation of a crystal. More

recently Jiang et al. [33] and other groups have found similar precursor

activity with other polymers examined with Fourier transform infrared

spectroscopy (FTIR) in situ during crystallisation and there are some

parallels in the recent work of Rabani, Reichman, Geissler and Brus [34] on

the formation of nanoparticle structures during drying.

Welch and Muthukumar [26] suggest entropic barriers are involved initially,

that chains attach themselves to the growing crystal in line with the existing

chains and that lamellar thickening takes place at a later stage in a

cooperative fashion.

Wurm and Schick [35] heated poly(ε-caprolactone) and syndiotactic

poly(propylene) with small laser pulses and presented evidence towards a

model with crystallisation first taking place by a partially ordered metastable

structure in the melt that becomes progressively more ordered into a

lamellar crystal as it undergoes a stabilisation stage.

Doye and Frenkel [36] disagree with the Hoffman-Lauritzen theory as to how

it predicts lamellar thickness and attempt to improve aspects of the

Sadler-Gilmer approach which is based on entropic barriers.

There are a number of other theories that also try to overcome the

weaknesses of the Hoffman-Lauritzen theory that was a major step forwards

over forty years ago. It is the Olmsted et al. approach that we will later see

24

supported in one aspect of the research being presented in this thesis. That

aspect is the minor phase separation seen in certain cases during

crystallisation and later re-melting.

The picture of lamellar structure and formation presented by Keller,

Lauritzen and Hoffman and others is an ideal one. In practice the loops of

the chain folds can either be close ones or they can re-enter the lamella

some distance away as in a “telephone switchboard” model.

Often a group of polymer chains will cluster together to form a lamella but

some of the sections of some chains will also be incorporated in other

lamellae. The intervening sections will meander through the amorphous

region between the lamellae.

A high undercooling below the melting temperature, either by rapid dropping

of temperature to a desired isothermal crystallisation temperature or by fast

dynamic cooling, leads to strong driving forces for crystallisation. The short

crystallisation period, which results from the strong driving force, does not

allow as much time for polymer chains to reorganise themselves into

favourable configurations. A far from ideal structure is then locked in.

The kinetics of crystallisation is strongly affected by the cooling rate, as

mentioned above. The crystallisation temperature is lowered as the cooling

rate is increased. Usually the amount of material that becomes crystalline is

reduced and the speed of crystallisation increases [37] with faster cooling,

resulting in a highly amorphous solid if the material is quenched, for

example, into ice water or liquid nitrogen.

It has been explained above that the question of crystallisation kinetics is a

complex one, even in the situation where crystallisation takes place

isothermally. There are good reasons for also wanting to study

crystallisation in a non-isothermal or “dynamic” context. That is the way

crystallisation takes place in almost every production environment so an

understanding of the processes under conditions emulating real life is also

needed [38].

There is also another consideration when crystallising a range of blends

where there may be multiple crystallisations. The temperature(s) for

isothermal crystallisation have to be chosen specifically for each particular

blend [6] and a wrong choice will obscure the information being sought. It

25

would be difficult to obtain meaningful results from a comparison between a

number of blended materials where the melting temperatures of the

constituents differ markedly. This is particularly so when differing

compositions of any pair being blended could lead to differing crystallisation

temperature depressions.

A practical solution is using non-isothermal crystallisation so that

crystallisation takes place when the molecules are ready for crystallisation at

the cooling rate used. A faster cooling ramp leads to a greater undercooling

before the crystallisation takes place because of kinetic factors. The

crystallisation does actually take place at near isothermal conditions

because the self-generated temperature field from the latent heat of

crystallisation does tend to maintain the local temperature in a pseudo-

isothermal condition during the crystallisation. Some authors utilise this

self generated pseudo-isothermal crystallisation in their own manner to

achieve specific undercoolings that would otherwise be difficult to achieve

[39]. It is not an ideal situation from a theoretical perspective but does

provide a practical way of solving the conundrum in those cases.

1.2.4.5 Lamellar melting

The melting of lamellae of monodisperse polymers takes place over a greater

temperature range than for small organic molecules or metals. This is

because the long polymer chains have to reorganise themselves and

dissociate themselves into the melt from their places attached to the side

faces of the lamellae. That process takes time and is at least partly

sequential with one layer being removed before the next one can also escape

into the melt.

The process of melting (some) individual chains from a number of lamellae

has been detected with polymer chains partly disengaging themselves into

the melt and recrystallising those sections onto the lamellae [40]. This was

carried out by using quasi-isothermal TMDSC using a very small amplitude

of the temperature oscillations. The reason for the small oscillations in

temperature was to keep the chains partly tethered to the lamellae so that

there was no barrier to recrystallising back onto the same lamellae.

The energy barrier to re-nucleation just referred to is a contributor to the

substantial difference in melting and crystallisation temperatures found for

polymers. It is necessary to have a substantial undercooling of perhaps

26

10-30 0C before the formation of lamellae takes place. That is the case even

when there are nucleation sites present from nucleating agents that have

been added to the polymer, or because the original melting of crystalline

regions was incomplete. This is discussed below.

Different lengths of polymer chain have slightly different melting

temperatures with higher molecular weight polymers having higher melting

temperatures than lower molecular weight polymers. The influence of this is

much greater at low molecular weights and negligible at high molecular

weights where the predominant factors are lamellar thickness and crystal

perfection. For example, Smith and St.John Manley [41] point to quasi-

monodisperse fractions of polyethylene with Mw = 1000 having a melting

temperature of 105 0C, that rising to 121 0C for Mw = 2000 but only rising

further to 131 0C for a molecular weight of 20,000. The same relationship

applies for crystallisation. Polydisperse polymers, as found in the real world,

therefore have wider melting and crystallisation temperature ranges than

monodisperse polymers and far wider than for small organic molecules.

1.2.4.6 Melting behaviour of semicrystalline polymers

We have now looked at both completely amorphous polymers and the

melting and crystallisation of lamellae. Even polymers in a very highly

crystalline state have 10% or more of amorphous material incorporated

between the lamellae and many semicrystalline polymers are 50-80%

amorphous.

The glassy amorphous material will become rubbery at the glass transition

temperature as the temperature is raised. The viscosity of a semicrystalline

polymer will be higher than with a fully amorphous one because the lamellae

act as relatively inert platelets restricting motion. Eventually the

temperature reaches the lamellar melting temperature and the polymer

segments that had been in the lamellae peel off the lamellae, becoming

indistinguishable from those that had been in the amorphous part. There

will be some drop in overall viscosity at the melting temperature (Tm) to that

of the amorphous material at that temperature as the melting lamellae cease

to restrict molecular motion in general.

1.2.4.7 Spherulites

Spherulites and other larger scale structures are made up of lamellae. The

spherulites are lamellar structures that have grown, splayed out and twisted

27

to create spherical forms. They appear as Maltese crosses under crossed

polarisation illumination. They are particularly important in polymers

crystallised from the melt. As early as 1888 Lehmann had made the

conclusion that they formed from long crystals that had forked as they grew

and spread out to fill up the space until they impinged on other growing

spherulites or until growth had stopped. The early observations of “twisted

crystals” in the 1920s and 1930s were for relatively small natural

macromolecules but by the mid to late 1940s they were being recognised in

polyethylene. Bryant identified that long polymer chains could partake in

multiple lamellae within a spherulite. That connection of the chains between

lamellae means that there is a coupled growth front for the spherulite as a

whole with the stacks of lamellae having amorphous material in between.

Other structures of lamellae that are encountered are axialites and hedrites

where crystals are attached to a common axis. Axialites can occur in

crystallisation from the melt but will be observed differently depending on

the direction of the axis relative to the observation direction. There are also

fibrillar structures encountered with oriented growth under stress such as

polymer fibre formation.

1.2.4.8 Poorly and partially crystallising polymer types

Aromatic groups on the main polymer backbone will have difficulty in

crystallising into lamellae because of steric hindrance. Polymers with long

branches will encounter difficulties in having the chains folding side by side

in lamellae because the side chains will get in the way sterically. Some

branched types will not crystallise in the main backbone but long side

branches may form lamellae. Copolymers often have one section that

crystallises and another part that does not.

1.2.4.9 Polymer-polymer miscibility

Often we want to utilise the strong points of two polymers to produce a

better material. The problem is that different types of polymers are usually

immiscible. There are a number of ways this can be overcome. One is to

synthesise block copolymers that have long chain segments of each polymer.

The synthesis in production quantities is often expensive. A number of

alternative approaches are possible, such as to use bulk quantities of each

polymer with a smaller amount of the relatively expensive copolymer to tie

phase separated domains of each type together as a compatibliser. Another

28

is to use maleic anhydride to reactively compatibilise the two. An alternative

method is to have two materials that can hydrogen bond together and use

this to compatibilise the materials as described by Huang et al. [42].

1.2.4.10 Polymer-diluent systems

The Flory-Huggins theory is an attempt to determine the ∆Gmix for polymer

solutions. Flory [43] and Huggins [44-46] independently put forward a

theory that has been modified by others in a variety of ways. The approach

is an extension of an earlier one by van Laar in which he treated two types (1

and 2) of equally sized molecules in an ideal solution as occupying the sites

of a three dimensional lattice. He then predicted the ∆Gmix as a function of

the universal gas constant, absolute temperature, numbers of moles of each

and the mole fractions of each material. That approach (which can be used

to derive Raoult’s law) was a failure for the case of polymer solutions. It was

extended by Flory and Huggins with the restraint that the segments of

polymer molecules within the solution are interconnected. It utilised an

interaction parameter Χ12 between the polymer and solvent molecules and

led to the ability to predict melting temperature depressions and phase

diagrams.

Flory [47 p.569] shows that

( )21121

10

11 vXvVH

RVTT u

u

mm

−∆

=− Eqn. (1-2 )

where mT is the equilibrium melting temperature of the mixture, 0mT the

equilibrium melting temperature of the pure polymer, R the Universal Gas

Constant, v1 the volume fraction of the diluent, ∆Hu is the heat of fusion of

the repeat unit, V1 and V2 the molar volumes of the diluent and unit

respectively and X1, the interaction parameter. The assumptions made in

the theory are that there is no volume change upon mixing, interactions of

the different types of segments cause the enthalpy of mixing after same type

interactions have been replaced, polymer repeat units and solvent molecules

are the same size and the number of combinations of polymer configuration

solely determines the entropy of mixing.

The Flory-Huggins lattice model, mean field theory above and its large

number of variants is suited to describing melting point depressions,

plasticisation and liquid-liquid (L-L) phase separations but not for

29

liquid-solid (L-S) phase transitions as taking place during crystallisation, as

discussed by Hu, Frenkel and Mathot [49]. No models to derive the

diluent-polymer interaction parameter Χ12 from linear relationships between

crystallisation temperature depressions and concentration have been located

in the literature. A number of the materials combinations studied in the

thesis do exhibit linear relationships between Tc and concentration for both

polyamide and for diluent. Calculations based on the above equation are

carried out in those chapters where melting depressions were found to occur

with a linear relation between melting point and concentration found.

Kelley and Bueche [50] use the additivity of free volume for the pure

materials in a miscible blend to determine the glass transition temperature

of a blend. Their calculations lead to good predictions of the Tg versus

composition curve and to the so-called Kelley-Bueche line delineating

vitrified and unvitrified blend regions. The intersection of the vitrification

and cloudpoint curves (due to phase separation) is the Berghmans Point.

1.2.5 Linear polyamides (Nylons) 1.2.5.1 History of polyamides

A linear polyamide was polymerised inadvertently at the end of last century

but was not recognised as being a polymer. A gelatinous mass had resulted

from experiments with amino carboxylic acids.

Prior to the 1920’s organic chemists failed to recognise the importance of

polymeric materials, concentrating their efforts on producing monomolecular

weight compounds. During the 1920’s, Staudinger recognised the existence

of polymeric material by relating solution viscosity to molecular weight.

Wallace H. Carothers was a brilliant organic chemist and in 1928 was

employed by DuPont to carry out research. He elected to continue in the

polymeric field opened up by Staudinger.

1929 was a period where great controversy still existed as to whether

polymers were long chain molecules, colloids, or aggregates of cyclic

compounds. At about this time Carothers [51] wrote a short review that for

the first time clearly identified the two main polymerisation reactions that we

now know as “chain growth” (addition) polymerisation and “step growth”

(condensation) polymerisation. He identified, in this review, that molecules

containing an amino group and a carboxylic acid group could condense to

form polymers such as polyamide-6. He also suggested that it may be

30

possible to condense diamine compounds with dicarboxylic acid compounds

to form polymers such as polyamide-6,6. In the early 1930s, when linear

polyamide-6 was being synthesised with caprolactam, he proceeded from

polycondensation with ε-aminocaproic acid to the synthesis using

hexamethyl diamine and adipic acid.

By 1939, the US approach had developed to the extent that there were

plants set up to commercially produce polyamide-6, which by then had

acquired the commercial name Nylon. This commercial activity was soon

subsumed by the war efforts. Germany pursued the investigation of optimal

polyamide types using the amino acid condensation and did not produce

polyamides commercially until after the war.

Nylons were one of the early polymers developed commercially. Nowadays,

they are manufactured industrially for a broad range of applications such as

clothing, stockings, carpets, fishing lines, tyre reinforcers, seat belts, and in

the components of a wide range of appliances and equipment. The fibre

component alone of linear polyamide worldwide production is in the order of

4 million tonnes per annum. This comprises nearly a quarter of total

synthetic fibre production, as noted by Elias [52].

Later developments lead to polyamides made with aromatic groups in the

main chain (called Aramids), branched polyamides and copolymers

incorporating polyamides in various forms. These later types are not covered

in this research work and they represent a much smaller volume of

commercial production.

1.2.5.2 Strengths

Polyamides are tough, impact resistant, flexible, abrasion resistant, heat

stable materials [53, 54] whose characteristic physical properties are mainly

determined as a result of hydrogen bonding. There are a range of

polyamides with varying properties dependant upon molecular structure of

the monomer repeat units.

Some of the newer polyamides such as polyamide-4,6 [55] have very high

melting temperatures, and mechanical stability that allow them to be used in

automotive applications near the engine. This particular polyamide has the

fast crystallisation that makes it attractive for injection moulders.

31

1.2.5.3 Weaknesses

Humidity plasticises and weakens polyamides. Polyamides also become

brittle when dry. Both these characteristics result from hydrogen bonding.

1.2.5.4 Chemical structure and polyamide types

Linear polyamides have a main chain with repeated amide units

incorporating -CONH- sections as shown in Figure 1-12. The amide unit is

always trans across the polymer backbone although it can sometimes be

partly twisted.

O II

– C – C – N – C – C - I

H Figure 1-12 CONH amide units found in polyamides showing the trans configuration of the bonds.

There are two basic types of linear polyamides, the polyamide-n type and the

polyamide-m,n type where the m and n are numbers representing the

number of carbon atoms in (parts of) the polymer repeat units.

Polyamide-n types, with n carbon atoms per repeat unit, can be formed by

condensation from amino acids such as in Carothers’ earlier work. Only one

material is used as the monomeric substance. An example is the ring

opening of caprolactam with its 6 carbon atoms and one nitrogen atom in a

ring. The opened ring is polymerised end to end into long chains forming

polyamide-6. Water is a by-product of the high temperature polymerisation

reaction and is pumped away to drive the reaction forward.

The polyamide-m,n types are obtained by the polycondensation of a diamine

and a dicarboxylic acid (or diacid). The number of carbon atoms in the main

chain due to the diamine gives the first number, m, and the number of

carbon atoms in the diacid gives the second number, n. For example,

hexamethyl diamine and adipic acid are used to synthesise polyamide-6,6.

Sometimes the number is placed before the word polyamide and sometimes

PA or Nylon is used. In some situations the name of the amino acid is used.

It is common to see PA-6, PA6, Nylon-6, Nylon6, 6-Nylon, Polyamide 6 and

poly(ε-caprolactam) for polyamide-6. There is an even greater variation of

naming for the m,n (or mn) types. Sometimes the comma is left out with

Nylon 612 meaning polyamide-6,12. The maximum length of polyamide-n

32

types is in the twenties and the maximum length of a polyamide-m,n is

similar so there is usually little confusion in omitting the comma.

When the numbers n or m+n are small then the repeat distance is shorter.

These polyamides are often referred to as “short” or “lower” polyamides as

distinct from “higher” polyamides. This does not refer to the number of

repeat units in the total polymer chain length.

It should be pointed out that polyamide-6 and polyamide-6,6 are quite

different materials even though the density of amide bonds in a polymer

chain is the same. The melting temperature of polyamide-6 at 225 0C is

some 30 0C less than for polyamide-6,6. The reasons for this will become

evident later.

The simplified structure of polyamide-6 and polyamide-6,6 are shown below in Figure 1-13 with repeat units in bold font. O O II II

- C - C - C - C - C - C - N - C - C - C - C - C - C - N - I I H H

polyamide-6 O O II II

- N - C - C - C - C - C - C - N - C - C - C - C - C - C – I I H H

polyamide-6,6

Figure 1-13 Repeat units of the n type polyamide-6 with all amide groups in the same direction and m,n type polyamide-6,6 from diamine combined with diacid and having the amide groups in alternating directions.

Note that the amide group is asymmetric so that the polyamide-6 repeat unit

fits head to tail along the molecule. The molecule as a whole is

unidirectional. On the other hand, the polyamide-6,6 can be seen to have

points of symmetry at the mid points of the amide and of the diacid groups.

This is an important point and will be taken up later. The “directional”

polyamide-n types will have antiparallel sections of molecules next to each

other as the molecule loops back in a hairpin. This happens as the

backward and forward laying of the molecule into place occurs on the lateral

33

faces of the lamellae. Sections of different molecules layered against them at

later stages can be parallel or antiparallel in direction.

It can also be seen that the polyamide-6 has a repeat length of 7 atoms in

the backbone whereas the polyamide-6,6 has a repeat length of 14 atoms.

Polyamide-6 and -6,6 are used for textiles because of their high tensile

strength. Polyamide-6,10 and polyamide-11 have longer distance between

amide groups are used for sutures and sporting goods requiring flexibility.

1.2.5.5 Biological-polyamide parallels

The biological fields touched on below are examples of, perhaps, the most

exciting potential areas to which this research could contribute because the

boundaries between biology and synthetic chemistry are breaking down and

both disciplines are learning from each other. This can be seen in a recent

review with over 160 references by Cunliffe, Pennadam and Alexander [56].

Linear polyamides are one of the most important natural polymers and are

known by biochemists and biologists as proteins or polypeptides. The

peptide linkage referred to by biologists is identical to the amide linkage that

occurs in synthetic linear polyamides. The molecular structure of

polyamide-2 forms a very simple model [57, 58] for a protein. Some of the

parallels between polyamides and proteins can be pointed out. A better

comprehension of polyamide crystallinity in different environments could

potentially lead to improved understanding of the way in which proteins fold,

recognised nowadays as a very important area of biology. Proteins can form

“molten globules” before crystallising out fully [32, 58] and this concept may,

in turn, be relevant to the way in which polyamides crystallise from the melt

or solution, particularly in the light of the recent work of Olmsted et al. [30]

on the formation of crystallites in molten polymers.

1.2.5.6 Polyamide Hydrogen Bonding

Polyamide crystallisation is more complicated than with many polymers

because hydrogen bonding constrains the crystallographic possibilities

further than just the steric considerations [59].

Hydrogen bonding, in general, was discussed earlier. It is now appropriate

to look at hydrogen bonding specifically in polyamides. The nitrogen atoms

in the amide sections are highly electronegative, withdrawing some of the

charge from the attached hydrogen. Normally the oxygen from the carbonyl

34

bond in another amide group elsewhere in the polymer chain or from

another molecule will be attracted to the hydrogen to form the N-H.…O

hydrogen bond, as portrayed in Figure 1-14.

O II

– C – C – N – C – I H . . .

O II – C – C – N – C – I H

Figure 1-14 Amide to amide hydrogen bonding found in polyamides showing the bridging from the electronegative oxygen of one amide group to the electron deficient hydrogen attached to the electronegative nitrogen atom of another amide group in the same polymer chain or another molecule

In general, there can be weak and strong hydrogen bonds. Those involved in

polyamides are considered moderate to strong.

These hydrogen bonds in polyamides are pervasive, being substantially

consummated in the amorphous state and are even present at a significant

level in the melt [60-62]. This makes the polyamides much more viscous in

the 50 0C range above their melting temperature than many other polymers.

They are the driving force that locks the crystallising lamella into one or

another crystalline form. They are also the reason for the very high melting

temperature of linear polyamides because they provide stability to the

lamellar structures.

Other molecules can be incorporated into the amorphous polyamide

structure, such as water, which plasticises and weakens polyamides by

displacing the hydrogen bonds.

1.2.5.7 Polyamide Crystallinity

Linear polyamide crystallinity is strongly affected by the exact type of linear

polyamide because of the limited combinations of the way hydrogen bonds

can be consummated within the constraints of the number of molecules

between amide groups. The orientation of the non-symmetric amide groups

in the chains also plays a strong role such as in the difference of 30 0C in

melting temperatures of polyamide-6 and polyamide-6,6 referred to above.

There are also steric limitations between the sections of molecular chains

35

lying next to each other and between different molecules in a lamella. These

differences between polyamides can be exploited to gain a better

understanding of polyamide crystallinity and the part hydrogen bonding

plays in their properties.

1.2.5.8 Polyamide Crystalline Structures

This and most of the following few sections are included mainly to provide

background understanding of the crystallographic forms polyamides can

take in differing situations and on Brill transitions rather than raise

expectations of the discussion of those in the experimental results.

We will first describe the five major crystallographic forms encountered with

polyamides as they crystallise from solution or the melt.

There are:

a) α, where the hydrogen bonds are in planes or “sheets” parallel to the edge

faces of lamellae (often intra-molecular bonds) and layers are built up

layer (sheet) upon layer [63]. With α there is an offset from the bonds of

one layer to another resulting in the basal planes of the lamellae being

inclined to the chain direction. Wide Angle X-ray Diffraction (WAXD)

gives two peaks at approximately 0.44 and 0.38 nm respectively at room

temperature. This is a stable crystalline form.

b) β, identical to α except that the chains with their offsets are stacked one

up and one down resulting in the (rougher at a molecular level) lamellar

basal plane being more or less perpendicular to the chain direction [63].

This is a stable crystalline form.

c) γ or pseudohexagonal and has inter-molecular hydrogen bonds between

amide groups in separate layers (sheets). The energetics result in a slight

offset between chains of different layers [63]. The chain spacing is nearly

hexagonal with a spacing of approximately 0.41 nm. With equal intra-

and inter-sheet distances between chains it is possible now for the

hydrogen bonds to be inter-sheet rather than intra-sheet.

d) Hydrogen bonds with more than one direction. Here, the amide groups

are twisted to give optimal energetics with one hydrogen bond to an

amide group in a chain in the same sheet and the next hydrogen bond

above or below being to the next sheet. Recently, polyamide-6,9 was

36

found by Franco et al. [64] to belong to this overall group of polyamides.

The groups that initiated this understanding are Subirana, Puiggali,

Navarro and colleagues with collaboration from Atkins, Hill, Cooper and

Jones [65-71].

e) Metastable pseudohexagonal forms [72] (broader single peak with X-rays)

and other forms with imperfect α structures.

There are also some other minor crystalline forms various authors have

referred to, including the Atkins, Hill, Hong, Keller and Organ [73] work

showing polyamide-4,6 has an α-like structure but with the chain direction

completely perpendicular to the basal plane and amide groups in the chain

fold. This is unlike the usual inclination to the basal plane, as found with

polyamide-6 and polyamide-6,12.

1.2.5.9 Effect of polyamide Type and Segment Length on Crystal Form

The exact way that a crystallising polyamide molecule folds backwards and

forwards to match up hydrogen bond acceptors and donors is very important

[48 Section 1.3]. We know from earlier work by Roberts and Jenekhe [74]

that virtually 100% of hydrogen bonds are consummated in the crystal, even

if it requires bending of bonds or the backbone of the chain to link through

from N-H to O. Both the parallel and antiparallel chain alignments can

connect hydrogen bonds easily within the molecule if the polyamide is an

“odd” numbered polyamide-n such as polyamide-7. Odd Nylon n types tend

to be more stable in the α- or β-form.

The stable form for “even” polyamide-n types, such as polyamide-6 is

generally the α− or β-form with hydrogen bonds matching up parallel to each

other and perpendicular to the overall polymer backbone. The angle in the

lamellar basal plane between the intra-molecular hydrogen bonds and the

corresponding chains of the next layer is at 67.50 to satisfy the steric and

energy constraints. Distances between chains within the sheets are greater

than between the van der Waals bonded sheets. The coefficient of thermal

expansion is less in the plane of the molecules than the inter-planar

direction. The hydrogen bonding constrains the molecular chains in a sheet

much more than between molecular sheets, as these are only held together

by the weaker van der Waals forces. The longer even polyamides can be more

stable in the γ form.

37

Even polyamide-n types in the α- or β-form have higher melting

temperatures than similar repeat length odd polyamide-n types by nearly

20 0C. The γ-form is regarded as being thermodynamically more unstable,

which correlates with the lower Tm.

The situation is further complicated with polyamide-m,n types because there

can be odd with odd, even with even, odd with even and even with odd

numbers of carbon atoms in the diamine and diacid sections respectively.

Even the last two are different in the way the parts of a molecule or parts of

different molecules can link together to consummate the hydrogen bonds.

The requirements for crystallinity are that this all happens in a consistent

way over (at least) regions of lamellae. In some cases the crystal repeat

distances are two monomer repeat lengths.

Odd-odd, odd-even and even-odd polyamide m,n types are usually more

stable in the γ-form. Here the hydrogen bonds are made between amide

groups in adjacent molecular sheets. The γ configuration has the two

hydrogen bonds in a molecular repeat unit at an angle to each other, and

neither is exactly perpendicular to the zigzagging backbone. This is because

the bonds do not exactly match up opposite to each other. The energy of the

total configuration must be minimised. It leads to hydrogen bonds holding

the sections of molecules further apart than would be the case for the α- or

β-form, and even further apart than for polyethylene. It also leads to a

slightly shorter repeat length. The total outcome is a crystal with slightly

lower density. The angle in the basal plane between the intramolecular

plane and the chains of the next layer is close to 600 and this leads to a near

hexagonal crystal structure, usually referred to as pseudohexagonal. The

coefficient of thermal expansion is the same in both directions of the basal

plane.

Even-odd polyamide-6,7 would seem, at first sight, to be similar to the

odd-even polyamide 7,6 but the bonds have to be twisted at different angles

to make the O.…H-N connections. The result is slightly different material

properties between the two polyamide types.

There are a number of diverse characteristics that can be found in the

different polyamide types. Even-even polyamide m,n types are generally

more stable in the α- or β-form [75]. “Shorter” Nylons have a higher

38

hydrogen bond density and have higher melting temperatures and densities

than the same types with longer overall repeat lengths. Polyamide-6,6 has a

30 0C higher melting temperature than Nylon-6, although both have the

same overall hydrogen bond density. Crystalline forms are different because

of the different orientation of amide groups within the chains [76, 77].

Different Nylon types have differing levels of moisture uptake due to their

various hydrogen bonding configurations [78 p. 324].

1.2.5.10 Multiple crystalline forms are possible - Polymorphism

There is usually more than one form possible for a particular polyamide type

but it often depends upon the thermal history as to which one is present in a

sample. The α- or β-form is more stable for longer polyamide n types and

the γ-form for shorter polyamide n types. Polyamide-6 appears to be equally

likely to have both forms, and these can coexist in a lamella. Polyamide-4,6

can exist in both α- and β-forms at the same time [79].

The form that exists in a polyamide depends on steric restrictions and the

most energetically favourable situation at a particular time. Often a

metastable crystalline configuration will form first, and later the crystalline

structure will change to another arrangement of hydrogen bonds, bond

angles and crystal cell distances. It can become energetically more

favourable to change to a different configuration as the thermal history of a

crystal develops.

Conversion between the two forms can be made to take place by temperature

changes [80] and also by solvents or materials that make polyamide swell

[81].

Sometimes a number of crystal forms will be present in the one sample [82]

and for polyamide-12 [83], the crystal structure can be varied by pressure

and cooling rates.

1.2.5.11 Effect of pressure on crystallinity, melting temperature and crystal form

Pressure often affects polymers by increasing crystallinity [84],

melting/crystallisation temperature and can change the crystallographic

structure. In particular, pressure affects the way in which polyamides

crystallise such as with the Ramesh work on polyamide-12 [83] and

supported by the English abstracts of the Chinese language work by Lu

Huang, Fan, Cai and Xie on polyamide-6 [85, 86]. Gogolewski and Pennings

39

show in their work on polyamide-6 [87, 88] that crystallisation under

pressure increases the crystallinity, although a greater increase can be

gained afterwards by annealing under pressure.

1.2.5.12 Metastability

Some materials go to metastable forms above or at the crystallisation

temperature and then change to more stable forms as the temperature is

lowered. Fast cooling can often trap crystal structures in a metastable form

because the molecules quickly lose the energy to surmount an activation

energy barrier. “Cold crystallisation” can often only take place when

previously quickly cooled material is raised in temperature to near melting.

A kinetic event takes place rather than a thermodynamic one with the

melting and recrystallisation into a more stable form before melting of the

stable form into liquid melt can take place.

1.2.5.13 Brill Temperature

The Brill transition occurs where a low temperature α form is heated so far

that the inter-sheet spacing increases until it is the same as the intra-sheet

spacing. Some contraction of the intra-sheet spacing is required with

temperature increase for the energy minimisation of the structure.

Eventually both d-spacings become 0.41 nm in a hexagonal structure. At

this stage, the hydrogen bonds can easily change from intra-sheet to

inter-sheet. The structure then becomes the γ form described above. The

changes in d-spacings between crystalline planes can be followed with

WAXD as temperature increases. The Brill temperature (TB) is the point

where there is no difference in spacings. The Brill transition has been most

extensively studied in polyamide-6,6 [64, 89-92] but does also occur in other

even-even polyamides [79, 93-98]. The Brill transition is reversible and on

cooling, the stable hexagonal γ form material reverts to α form.

Some polyamides do not quite reach the Brill transition before they melt.

The stable form of the crystal will remain α- or β-form in these cases.

Kohan states [99 p. 143] that Brill transitions are usually not seen with DSC

scans for melt crystallised samples, so they are not expected to be seen with

DSC in the work carried out here.

40

This discussion has been included to alert to some of the complexities

involved. No further discussion of Brill temperatures is given in the text

because of the use of DSC results and absence of WAXD results.

1.3 Relevant papers in the area to be covered in the research There has been much done in the way of research on amorphous-amorphous

and crystallisable-amorphous polymer systems (including polyamides) and

methods to overcome miscibility problems. Much of that has been driven by

the desire to improve the physical properties of polymers in a cost-effective

manner. A few have done work on semicrystalline-semicrystalline blends,

sometimes enhancing miscibility by hydrogen bond interactions [100]

(although Qiu et al. only touch on those interactions) and a few have

researched semicrystalline-(crystalline) small molecule blends.

The area covered by the research is concerned with the melting and

crystallisation of aliphatic polyamides with certain, potentially hydrogen

bond disrupting, small molecules. With the exception of water, this area

does not appear to have been covered by other researchers but there have

been papers published in adjoining areas and these will be reviewed in this

section.

1.3.1 Small molecule-small molecule Sucrose is usually crystallised from anhydrous melts or highly concentrated

solutions in a controlled manner to generate the specific textures or

appearance required for the final product such as fudge, hard candies. It is

shown in a paper by Bhandari and Hartel [101] by DSC and XRD results

that it is possible to reduce the crystallinity from molten anhydrous sucrose

to about a third by the addition of up to total weight 20% fructose, glucose

or a mixture of the two in a co-crystallisation process.

1.3.2 Polymers with small molecules Kristiansen et al. have studied sorbitol based nucleating agents(DMDBS) for

removing haze from isotactic poly(propylene) (iPP) in their recent paper

[102]. They also studied a very much wider range of concentrations than the

optimal clarifying concentration near 0.8% so that they could understand

some of the mechanisms. They refer to a regime III near the melting

temperature of DMDBS (higher than the iPP) where phase separation takes

place (as determined with optical microscopy). At lower temperatures, there

41

is a partially crystallised fibrillar structure that solidifies below the eutectic

temperature.

Simek et al. [103] have studied the melting temperature depression of

isotactic poly(propylene) (iPP) by alkanes. They have used the Flory Huggins

relationship to explore the size of the effect.

Kim and Kim [104] looked at liquid-liquid phase separation occurring with

vinyl acetate and paraffin wax blends with poly(ethylene-ran-vinyl acetate)

using DSC, cloud point determinations and wide angle X-ray diffraction.

Their 1 0C/min and 10 0C/min cooling results are closest to the 2 0C/min

and 25 0C/min cooling rates used in this work and show only slight

differences in the DSC thermograms for reheating after nonisothermal

crystallisation.

1.3.3 Blend interactions and hydrogen bonding There are some similarities in a recent paper by Rocco et al. [105] to the

original concept for hydrogen bond disruption by small molecules. In their

case they were interested in suppressing crystallisation of poly(ethylene

oxide) (PEO) by hydrogen bonding poly(bisphenol A-co-epichlorohydrin)

(PBE) to enhance properties of PEO being used as polyelectrolytes in

batteries. The blend interactions were observed with a shift for O-H from

3495 to 3348 cm-1 in going from the “free” (non-hydrogen bonded) state to

the “bound” hydrogen bonded state. This change of 50 cm-1 seen in peak

position (without Gaussian deconvolution) is indicative of what could be

expected with polyamides and hydrogen bond disrupting diluents if there

were any hydrogen bond interactions.

Dormidontova and ten Brinke [106] tackle the influences of hydrogen

bonding on micro- and macro-phase separation from a theoretical

perspective for comb copolymers with hydrogen bond interacting

end-functionalised oligomers.

Kobori et al. [107] looked at interfacial interactions of immiscible polymer

blends (linear-low density polyethylene/poly(methyl methacrylate) with

polyethylene) where hydrogen bonding did and did not play a role. The two

combinations respectively linear-low density polyethylene/poly(4-vinyl

phenol) containing polyethylene-block-poly(methyl methacrylate)

(LLDPE/PVPh with PE-b-PMMA) and the non-associating blend

42

LLDPE/PMMA with PE-b-PMMA were studied. FTIR measurements showed

differences with an extra peak for hydrogen bonded carbonyl groups 30 cm-1

from the normal peak for unbonded carbonyl groups at 1730 cm-1. These

were associated with differences in the phase boundaries demonstrating

lower interfacial tension between the phases.

1.3.4 Polyamides and Polymers Much of the work reported in this area is by researchers trying to overcome

the abysmal performance of polyamide/other polymer blends using a variety

of compatibilisers. Often comments are made about the uncompatibilised

blends that give an idea of the normal situation.

Moon, Ryoo and Park [108] discuss their work on using maleic anhydride

grafted polypropylene as a compatibiliser to improve

polyamide/polypropylene blends that are a semicrystalline/semicrystalline

combination.

Jafari et al. [109] studied the crystallisation of polyamide-6/polypropylene

blends using hot stage microscopy to look at the formation of polyamide

spherulites and how the polypropylene crystallised at a later stage. It will be

raised in a later chapter later that this paper may be relevant to the way

polyamides crystallise in certain circumstances.

Murthy et al. [110] took the interesting approach of blending a

non-crystallisable aromatic polyamide with (normally crystallisable)

polyamide-6 and used simultaneous small and wide-angle X-ray studies to

probe the crystallinity of the polyamide-6 in the blends. They found that the

polyamide-6 crystallinity was depressed by the presence of this other

polyamide.

Kim, Cho and Yoon [5] have recently studied the effects of compatibilisers on

blends between polyamide-6 and poly(vinylidenefluoride) (PVDF) to improve

the poor mechanical performance of these semicrystalline/semicrystalline

blends. The uncompatibilised blends had strong phase separation. The two

areas the uncompatibilised blends were noted for were poorer compatibility

in the amorphous regions and faster crystallisation.

PVDF is crystallised isothermally with polyamide-11 by Li and Kaito [111]

and studied as uniaxially stretched films with SAXS and WAXD with or

without annealing. There are limited DSC results with a peak for the

43

polyamide crystallising in the blend at a temperature higher than normal for

polyamide. It is possible that this experimental result is consistent in

mechanism with a couple of similar examples of this in this work, despite

theirs being a polymer-polymer system.

1.3.5 Polyamides and small molecules Cha et al. [8] have studied a system of polyamide-12 with poly(ethylene

glycol) (PEG) of differing molecular weights in relation to the formation of

membranes by thermally induced phase separation. This is chemically the

closest of the available literature to the systems studied here but they have

tackled it from a different perspective with a focus on the effects of molecular

weight of the diluent and on diluent-rich domain size. Their work used light

transmission changes to detect phase separation in dynamically cooled

(10 0C/min) melt blends. Samples 200 µm thick between coverslips could be

cooled quickly into the unstable region where droplets of polymer-poor

material formed and solidified once the phase separation temperatures had

been determined. Sample thickness at 200 µm is less than half the

estimated thickness of the samples investigated in this thesis but that is not

expected to induce significant differences due to dimensional constraints.

Some samples were initially produced by first solvent casting (with heated

vacuum drying) before forming the melted film. The authors claim that no

differences were detected due to this procedure. Videomicroscopy was also

used for phase separation temperatures. Their study gives experimental

phase diagrams with cloud point curves delineating the liquid-liquid phase

separation boundary and melting & crystallisation points for the Nylon at

different polyamide concentrations. They do not describe how these latter

data points were determined and whether they are the melting and

crystallisation peak temperatures or onset temperatures. However part of

the group describe in a later article how they use optical methods to

measure melting and crystallisation temperatures [10]. Whether those are

the methods used in the 1995 paper is not clear but that paper does not

describe the use of equipment other than hot stage optical microscopy

observations, videomicroscopy and SEM. The early part of the paper is

devoted to the development of phase diagrams giving the

temperature-composition conditions where two-phase behaviour exists and

the results of that are used for setting up experiments where mixtures are

44

quenched to 170 0C, diluent rich domains develop and later the samples are

cooled to ambient temperatures. The focus here is on the size of PEG-rich

domains that can be used to form membrane pores.

Their study covered PEG with molecular weights of 200, 400, 600, 1000,

1540 and 3400 Daltons. The results for the PEG having a molecular weight

of 200 Dalton are the closest to the carbazole and phenothiazine used in this

study (167 and 199 Daltons respectively) and PEG has a quite different

molecular form to the poly-aromatic rings of the carbazole and

phenothiazine. The carbazole and phenothiazine used here have quite

different molecular shapes to the PEG and only have molecular weights near

200. The molecular weight and form factors will affect the mobility of the

diluent molecules in the amorphous and molten polyamide in different ways

and the chemical potentials of the diluents with respect to that of the

polyamides will differ. Polyamide-12 is a polyamide with a lower density of

amide groups than the polyamide-6,12. It has a lower melting temperature

than the polyamides studied here and would be expected to have a lower

crystallinity also, approximating a polyolefin much more than them.

The major findings of the paper were that UCST behaviour was seen,

solid-liquid as well as liquid-liquid phase separation were seen, that the two

phase region was larger with increasing diluent molecular weight. They also

showed that both diluent molecular weight and content in the mixture

affected the interaction energies derived with Flory-Huggins theory. Factors

they found that are perhaps of lesser interest in the context of this thesis are

that the domain size was larger for greater PEG molecular weight but this

was not so strong an effect at low PEG content of the mixture.

This thesis later discusses the relative crystallisation temperatures of

polyamide and carbazole or phenothiazine diluent. It can be noted at this

point that PEG200 has a higher crystallisation temperature than

polyamide-12 in the study by Cha et al. and that this corresponds to the

case with polyamide-6, polyamide-6,9 or polyamide-6,12 combined with

carbazole.

45

1.4 The focus of the research project Aliphatic polyamides, or “Nylons”, are an important class of engineering

polymers. They are characterised by relatively high melting temperatures,

high impact strength and toughness.

An overall study of the literature in this general field has not uncovered

much work generally in the area of polymers melt blended with small organic

molecules and only one relevant paper on polyamides with small organic

molecules with that one looking at quite different aspects [8]. The literature

has shown that some others have achieved hydrogen bond complexing

between poly(ethyleneoxide) with poly(bisphenol A-co-epichlorohydrin)

whereas that was found not to occur with the materials chosen.

The research problem is to understand the processes involved in forming

high temperature solutions by melting linear polyamides with carbazole or

phenothiazine (as examples of small molecules) and in their crystallisation.

It is also to understand the resulting morphologies arising from

crystallisation and the effect of polyamide type.

The research had originally been planned to investigate the role of hydrogen

bond formation on crystallinity in linear polyamides. The concept was partly

based on the work of Damman, Point and coworkers [21, 112-114] in

creating molecular complexes between poly(ethylene oxide) (PEO) and

p-nitrophenol or resorcinol that are hydrogen bond complexed with them.

There was also some (as yet unpublished) work done by others within our

group at the University of South Australia on poly(ethylene glycol) (PEG) and

resorcinol.

The project had also been undertaken to study the effect of hydrogen bond

complexing on the physical properties that make Nylons desirable to use in

many applications. The potential benefits of the project were to aid in

widening the manufacturing/processing window for Nylons, to provide

options for adding dyes to Nylons in solution or the melt and the potential to

assist in developing new ways to deliver drugs within the body by

encapsulating them in polyamide excipients. A more fundamental reason for

doing this research was to help the understanding of hydrogen bonding in

polyamides in general. There was also the possibility of using a synthetic,

46

model compound to better understand protein folding because of the

similarities between amides and the peptides found in proteins.

The aim was to insert organic materials in the melt that would disrupt the

polyamide-polyamide hydrogen bonding so that the strong hydrogen bonds

would be destroyed and the material properties altered. This approach using

organic hydrogen bond disruptors is quite different from the earlier inorganic

approaches by others. Those had concentrated on iodination [115-119] or

the use of metallic ions such as Ca++ [120] for the study of changes in

crystalline structure. This alternative approach was taken because of the

obvious close parallels with many biological systems. The work was done in

the melt rather than room temperature solutions to avoid the three-way

competition for hydrogen bonds that would arise from dissolving the

polyamides in a solvent [42]. Polyamides require very strong solvents such

as formic acid, concentrated sulphuric acid, m-cresol or special solvents

such as 1,1,1,3,3,3hexafluoroisopropanol [121] that have to destroy the

polyamide-polyamide hydrogen bonds in order to dissolve the solid polymer

in the first place. The intention was to use DSC as part of the material

property analysis.

It will later be clear from FTIR results that the materials chosen did not

result in hydrogen bond interactions with the polyamides along the lines

expected [42, 112, 113].

1.4.1 Materials chosen 1.4.1.1 Polyamides

Four, quite different, representative polyamides were chosen for the study so

that the conclusions could be as general as possible. It has been shown

above that the melting temperatures and crystallography of polyamides are

influenced strongly by the type of polyamide. The relevant parameters

included whether they were polyamide-m,n or polyamide-n types, whether

polyamide-m,n is even-even or odd-even and the density of amide groups in

the backbone is high or low.

One of the polyamide materials chosen was polyamide-4,6 [122, 123] which

is an even-even polyamide with a high amide group density. It has much

higher crystallinity and melting temperature than most other polyamides

[55] due to the above factors. It provides one end of the scale of even-even

polyamides studied. Polyamide-6,12 is towards the extreme of even-even

47

polyamides having low amide density and still readily available. The very

common polyamide-6,6 was not chosen because there were no samples

available that were not known to have fire retardants and other additives

and because the two more extreme members of even-even polyamides were

being studied, allowing estimates for the intermediate polyamide-6,6.

Polyamide-6 was available in grades not known to have additives.

Polyamide-6 and polyamide-6,6 are the most common of commercial

polyamides so polyamide-6 is representative of both a mid amide density

polyamide and a polyamide-n type.

Polyamide-6,9 was also available in a grade not known to have additives and

is representative of an even-odd polyamide-m,n. Its melting temperature is

lower than polyamide-6,12, with even lower amide density, due to the

even-odd configuration having unfavourable hydrogen bond linkages. It is

also a member of the group of polyamides now known to have hydrogen

bonds in multiple directions [64].

These four polyamides have melting temperatures between 209 and 290 0C

and provide a compact group with a suitable range in repeat unit types,

stable hydrogen bond structures and melting temperatures. This should

allow us to draw some general conclusions about the interactions and phase

behaviour of polyamides melted with the two chosen materials.

1.4.1.2 Small molecules

Work started originally with 2-methyl resorcinal and p-dihydroxybenzene

(hydroquinone) as these had been hydrogen bond complexed with PEG in

Paternostre, Damman and Dosiere’s work [124, 125]. Evaporation was an

immediate problem because the polyamides melt at such high temperatures,

so other potential materials such as benzophenone with higher boiling

temperatures were also tested.

There were several determining factors in the choice of small hydrogen bond

disrupting molecules to be used in the originally planned melt complexing

project. A list of criteria was then drawn up.

Polyamides start to degrade (scission of the polymer chain at the amide

groups above approximately 325 0C as seen in the TGA thermogram later

(Figure 1-16 in Section.1.5.1). There is usually further polymerisation of

polyamides at temperatures near the melt and above [48]. Extended periods

48

at elevated temperatures above 300 0C would result in a marked increase in

polydispersity from scission and further polymerisation that would detract

from the validity of the work because of the uncontrolled molecular weight

distribution. The small molecule melting temperature upper limit became

300 0C.

a) Trials need to run substantially above the melting temperature of both

the Nylon and the potential disruptor so that self-seeding nuclei from

either material would not remain in the melt to cause premature

crystallisation. In particular, the polyamide should have over five

minutes fully in the melt to remove the previous crystalline state.

b) The potential disruptor should not evaporate or decompose at the

temperature of the trials, ie. more than 300 0C in the case of experiments

with polyamide-4,6. It was preferable to have the same material(s) for all

polyamides so that valid comparisons could be made.

c) The affinity of a hydrogen bond disruptor for the Nylon hydrogen bonds

should preferably be greater than the strength of polyamide-polyamide

hydrogen bonds.

d) There should only be one potential hydrogen bonding site on the molecule

so that bridging between several polyamide chains (or within a chain)

would be avoided.

These criteria are quite difficult to meet. For example the common Nylon

plasticisers N-ethyl o- or p-toluenesulfonamide boil at 196 0C and have

multiple potential hydrogen bonding sites per molecule. Another,

N-butylbenzenesulfonamide, boils at 314 0C

Nearly a dozen potential compounds were found that seemed to be suitable

and each one had a single N-H or C=O bond available for hydrogen bonding.

Some of those were not commercially available and could also not be

obtained via contacts in various laboratories. A handful of the rest

remained. Many of those were evaluated with Simultaneous DTA-TGA (SDT)

to eliminate poor performers on the critical evaporation criterion.

That left only two, carbazole and phenothiazine, that were reasonably

suitable. The melting temperature of carbazole is 246 0C and its boiling

temperature is 355 0C whilst the melting temperature of phenothiazine is

49

186 0C and its boiling temperature is 371 0C [126] . It was found that the

boiling temperature is not as critical as the vapour pressure at the working

temperatures near 310 0C. It will be seen in later chapters that carbazole

with the lower boiling temperature gave less trouble in the trials than the

higher boiling phenothiazine.

Problems were still encountered for carbazole and phenothiazine with

evaporation, even for polyamide-6,9, with its low melting temperature, and

even with high heating and cooling rates to minimise evaporation.

The structures of the two materials are shown below in Figure 1-15.

NH H

N

S Figure 1-15 The structures of Carbazole and Phenothiazine.

These are both relatively flat molecules although the phenothiazine has a

slight curvature from top to bottom as computed for us by Dr. Adam

McCluskey at Newcastle University in New South Wales, Australia. Both

have pi electron clouds above and below the benzene rings.

1.4.2 Sample blending and notation used for blends Small samples could be made up from powders in Differential Scanning

Calorimeter (DSC) pans to understand the initial melting (plus crystallisation

and later remelting) in the DSC. Larger blend samples were mandatory to

study properties using a variety of the techniques described below in Section

1.6 and these could be made in ampoules.

A consistent notation is used within the thesis for melt blend samples.

Polyamides are often described in the literature in various forms. For

example, polyamide-4,6 is seen in articles as Polyamide-4,6 polyamide-46

Polyamide4,6 polyamide46 PA-4,6 PA-46 PA46 Nylon-4,6 Nylon-46

Nylon4,6 Nylon46 4,6-Nylon and some other variants. The versions that

will be used here for melt blending are PA46, PA6, PA69, and PA612 for

polyamide-4,6, polyamide-6, polyamide-6,9 and polyamide-6,12 respectively

when combined with Car for carbazole, or PTh for phenothiazine (This

should perhaps have been PhTh but the aim was to keep it to three letters

signifying which diluent was involved in the blend.). Polyamide-4,6 blended

50

with carbazole is generally noted as PA46Car. Specific samples with known

percentages of polyamide are preceded by the weight percentage of

polyamide eg 39PA69PTh for a sample of 39% polyamide-6,9 in combination

with phenothiazine. The value of 39% would be calculated from the few

milligrams of each material used when blending in pans in the DSC or from

the TGA results where a sample is taken from the bulk material made in

larger quantities in an ampoule. The samples for TGA are taken from next to

the DSC samples. This notation provides an easily recognisable and

compact descriptor for each sample.

1.5 Experimental Techniques Used This section includes results that will be used to illustrate certain recurring

features that will be discussed throughout the thesis. The major focus of the

work rests on the results of Differential Scanning Calorimetry and Fourier

transform infrared spectroscopy (in Mid and Near infrared ranges) with the

support of Thermogravimetric Analysis for determining polyamide

concentration in ampoule samples.

1.5.1 Thermogravimetric Analysis Thermogravimetric Analysis (TGA) is used in this project to determine the

weight percentage of polyamide in a bulk sample where the composition may

vary markedly from the average for the whole sample. It is a technique

where a sample of material is heated in a gas stream with a furnace and the

weight is monitored accurately with an extremely sensitive balance.

Figure 1-16 Evaporation of carbazole followed by degradation of polyamide-4,6 in TGA.

51

This technique can be used because carbazole (or phenothiazine) in a blend

sample will evaporate in an inert gas stream (nitrogen) before the polyamide

begins to degrade. Evaporation usually takes place (at 10 0C/min ramping

rate) in the range 175-275 0C but the polyamide does not begin to degrade at

that ramp rate until well into the molten state over 325 0C. It means there is

a plateau in the TGA thermogram of remaining percentage of the samples’

weight vs. temperature at least in the range 275-325 0C. A small amount of

degradation products from the polyamide usually remains by 600 0C [127].

The plateau is clearly observable in the typical TGA thermogram depicted in

Figure 1-16.

1.5.2 Differential Scanning Calorimetry Differential Scanning Calorimetry (DSC) was the main technique used in the

experimental work. This is because it was able to provide information on

melting and crystallisation temperatures and crystallinity of samples when

they were being heated into the molten state and crystallised during cooling

to room temperature. Additionally, the DSC was used as a ‘furnace” to take

small samples of polyamide and diluent powders to the melt to study the

high temperature solutions. Monitoring could take place in situ whilst

carrying out this preparatory process. It enabled a better understanding of

the initial eutectic formation from the raw mixes of powders.

DSC measures the flow of heat into or out of samples when they are heated

or cooled. Thermal transitions as a function of temperature and time give

quantitative and qualitative information regarding physical (and chemical)

changes such as melting, crystallisation, recrystallisation glass transition

temperatures, cold crystallisations, polymerisation, degradation reactions,

volatilisation or changes in heat capacity. Melting and crystallisation

temperatures can be determined. The amount of crystalline material that

has melted or crystallised can be determined and, by comparison with

literature values for 100% crystalline material, the crystallinity can be found.

There are two types of instrument, “Heat Flow” and “Power Compensated”.

In the first type, the temperature difference between a reference and sample

pan is measured as both are heated in similar situations in a DSC cell. The

other type determines the amount of power required to keep the sample at

the same temperature as a reference as they are both heated in a cell.

52

There are a number of quite different designs for the cells used with both

types of instrument [128 p. 129]. The type used in this work is a Heat Flux

instrument. Figure 1-17presented here shows a cross-section of a DSC cell

for a TA Instruments calorimeter Model 2920 DSC.

Figure 1-17 Cross section of DSC cell (taken from [129 p. 4-5].

There are two methods of treating heating ramps for differential scanning

calorimetry. There is “standard” DSC with a constant ramping rate and

“temperature modulated” DSC (TMDSC) where a sinusoidal or sawtooth

[130] modulation is superimposed on the constant ramping rate. This later

method, developed since 1993 [131], was put forward as having a number of

experimental advantages. It has, however, been more recently recognised

that there can be limitations in the interpretations [132-134], especially with

melting and crystallisation events. The work in this thesis was done under

TMDSC conditions (with the extra calibration required) to utilise the smaller

sample size, increased resolution and sensitivity. Analysis could then still

be done at a Reversing/Non-Reversing level where it was required and

appropriate. Glass transition temperatures are also obtainable where the

crystallinity is not too high. Unfortunately, polyamides are often high in

crystallinity, leading to weak glass transitions.

There is a more extensive discussion of standard DSC, TMDSC and the use

of Lissajous figures to better understand thermal events during TMDSC in

Appendix B. The caveats placed on the use of TMDSC described in this

appendix mean that it was inappropriate to analyse the melting and

crystallisation processes of highly crystalline diluents and very crystalline

polyamides from a TMDSC perspective.

53

Small-molecule diluents remain solidified until the temperature is raised

sufficiently that molecular motion catastrophically breaks down the crystal

structure. The amorphous part of a pure polymer will be reduced in

viscosity with heating to the viscosity at melting and polymer chains

comprising the lamellae will “melt” into this fluid of the same composition.

Blends differ from both of these in that the amorphous part of a polymer is

highly plasticised by the diluent, forming a solution that is of lower viscosity

than the normal polymer melt. The lamellae essentially “dissolve” in this

liquid. Technically the correct usage throughout the thesis should be

dissolution but in many cases there are pure materials melting and blends

dissolving under the same section heading or in the same thermogram. The

common term “melting” has been used for both headings and figure captions

as well as text describing the melting of pure materials. Usage of

“dissolution” has generally only been followed in the text for blends where

there is specific discussion of polymer chains being removed from lamellae

into the liquid.

1.5.2.1 Thermogram Overlays

In general, the DSC thermograms are displayed as overlays with several

thermograms together in a figure. That is done to make better comparisons

between different compositions investigated under the same conditions. The

thermograms are all shown as heat flow in J/g against temperature in

degrees Celsius. All thermograms have exotherms pointing upwards.

The thermograms are spaced out vertically and coloured in a consistent

manner to aid clarity. The colour scheme can be seen in Figure 3-16 in

Section 3.4.4.1. The peak with largest amplitude starts with the endotherm

or exotherm near zero. In practice, that will be either carbazole or

phenothiazine. The other thermograms are placed in order of concentration

through to the polyamide so that trends can easily be seen as they relate to

polyamide concentration. It means that the polyamide curve will be at the

bottom for melting and at the top for crystallisation. The legends are always

with the pure polyamide (100% polyamide) at the top and range down to the

diluent (0% polyamide) at the bottom.

In some cases, the phenothiazine or carbazole peak is extremely large in

amplitude compared to the thermograms of the polyamide or the blends. In

54

those cases the very large peaks have been truncated in the figure so that

the detail of the other materials and/or combinations can be clearly seen.

First time thermograms of blends are generally drawn with a dash and the

repeat runs in the DSC are drawn with a solid line. The exceptions are the

few cases where there is more than one thermogram in the same

concentration range and other line types have been used.

1.5.2.2 Thermograms expected from thermal events

We will now consider the general forms of thermograms resulting from

different types of thermal events. This will facilitate discussion of results in

later chapters.

The DSC thermogram will have a single peak for melting or crystallisation if

the percentage of polyamide is exactly that for the eutectic composition

because at the eutectic triple point the solid changes at one time through

from the solid to the liquid phase or vice versa. The temperature of that

peak will be close to the equilibrium eutectic temperature but will be

modified by the dynamic heating or cooling not exactly being at equilibrium.

There will also be differences in heating and cooling eutectic peaks because

polymers are involved in this study and the normal melting and

crystallisation of polymers do not take place at exactly the same

temperature. Polydispersity of the commercial polymers used will also have

an influence on the outcome.

Consider now the case of polyamide/diluent with a polyamide concentration

different from the eutectic concentration and being heated. In the first stage

of heating, polyamide and the diluent melt up to the limit of solubility of one

material in the other in a eutectic melting peak. That peak temperature is

virtually constant across a wide range of total composition in samples.

There is now a residual of one or other material because the two materials

are not present at exactly the eutectic composition. The solubility of the

excess material will generally increase rapidly at higher temperatures. The

endothermic curve in the thermogram above the eutectic melting peak is due

to the progressive melting of more and more of the residual material as the

temperature is increased in the heating ramp. Eventually the excess is

consumed and the sample is completely liquid with no further melting

55

activity. This can be seen in Berghmans’ chapter of Mathot’s book [135

p. 214 Fig.8.7].

This process results in an endothermic curve above the eutectic melting

peak in the thermogram that takes the shape seen in the second peaks of

Figure 1-18. These peaks have a similar form due to a similar process

taking place, however, the second peak for 25PA6Car extends further as

more carbazole has to be dissolved into the liquid, requiring higher

temperatures to increase the solubility. A higher level of carbazole again

would require even higher temperatures to dissolve all the diluent. This

form of curve will be referred to in the text as a Temperature Limited

Solubility (or TLS) peak. It is interesting that the peak temperature is just a

few degrees before the end of the melting process that defines the totally

liquid state. A plot of melting peak temperatures against composition will be

seen in later chapters to take on the general form of the eutectic phase

diagram (Figure 1-6 of Section 1.2.1.6). There are, however, differences

because the peak temperatures are not the end of melting but peak melting

rate and because the system is not in an equilibrium state.

The forms of the curves are slightly different between excess of diluent and

for excess of polyamide but the principles are the same.

Figure 1-18 Two examples are shown to illustrate this general feature. The upper curve is for 25% polyamide-6 in carbazole and the lower curve for 64% polyamide-6 in carbazole. The first peaks near 195 0C corresponds to melting the eutectic composition and the second peak to melting the “excess” mixture of which there is more in the 25% sample.

56

We have seen in Section 1.2.5.12 that we can have metastable crystalline

forms locked in to lamellae, particularly by fast cooling. These metastable

lamellae have lower melting temperatures than the stable form and undergo

a melting and recrystallisation into the more stable form before the final

melting of the stable form. That can be observed in several variants. We can

see the first melting of the existing metastable crystals absorbing energy,

and later the heat given off in crystallisation of the metastable lamellae prior

to the main peak endothermic melting of the stable form for the

polyamide-6,12 sample in the thermogram below. We can also see a minor

version of these processes taking place for the 60PA612PTh thermogram in

Figure 1-19. This latter thermogram only shows a shoulder early in the

main melting peak.

Figure 1-19 Melting and recrystallisation of metastable crystals before melting the stable crystals. This can either be an extensive endotherm and exotherm pair, as with the pure polyamide, or there can be a subtle dip before the main peak and a shoulder on the leading edge of it for the blend.

1.5.2.3 Assignment of “Spiky” Crystallisations to Carbazole or Phenothiazine

The crystallisation of carbazole and phenothiazine take place extremely

rapidly because the molecules are quite small compared to long polymer

chains. The heat released in crystallising often makes the peak temperature

of crystallisation appear higher than the crystallisation onset temperature.

The form of the crystallisation peak is very distinctive, as can be seen in

Figure 1-20. It is very easy to identify a crystallisation as being from nearly

57

pure carbazole (or phenothiazine), unlike the situation during melting. The

following discussion about carbazole applies equally well to phenothiazine.

The distinctive slight rise in temperature is due to the sample thermocouple

being on the underside of the constantan dimple where the sample pan

rests. The pan contains the molten carbazole that is being cooled. The

carbazole is still molten at the time the thermocouple reduces in

temperature to below the carbazole freezing temperature due to slight

thermal lag in the system. This is because of small but noticeable thermal

resistances between thermocouple and carbazole. The freezing carbazole

within the sample maintains it at the carbazole crystallising temperature so

the thermocouple soon rises again to match that temperature. It takes

approximately 6 s for the heat flow to reach a maximum. In that time the

“Sample” temperature measurement “appears” to increase by 0.44 0C.

Figure 1-20 Displaying the radically different forms of crystallisation peaks for polyamide and diluents allowing identification of the material crystallising. Phenothiazine has a similar crystallisation thermogram form.

The crystallisation of polyamide, however, approximates a broader Gaussian

distribution because it is a polymer crystallising and because the polymer is

polydisperse (Section 1.2.4.5).

1.5.2.4 Phase diagrams derived from thermograms

Three examples are given here of experimental non-equilibrium phase

diagrams. Figure 1-21, is from Cha et al. [8] referred to earlier as closest to

the systems studied here. It used PEG (Mw = 200 Dalton) as the diluent with

58

polyamide-12 as the polyamide with cloud point measurements on

Liquid-Liquid phase separation-and some (undefined) melting and

crystallisation measurements.

Figure 1-21 Experimental phase diagrams measured under the condition of 10°C/min cooling rate: (a) nylon 12/PEG2 00 blend from Char et al. [8], where Tcloud is from cloudpoint measurements, and Tm & Tc are melting and crystallisation temperatures respectively.

130140150160170180190200210220230240250260

0 10 20 30 40 50 60 70 80 90 100Polyamide concentration (%wt)

Tem

pera

ture

(0 C)

TmPA69PureTmCarDeprTmEutTcPA69PureTcCarDeprTcEut

Liquid

Solid

Liquid & solid

SolidLiquid

Solid Liquid

LiquidSolid

Figure 1-22 Example of eutectic style non-equilibrium phase diagrams for heating to the liquid state and cooling raw materials and blends of polyamide-69 (PA69) and carbazole (Car) Melting peaks are noted with Tm and crystallisation by Tc. Eutectic points are denoted by TmEut or TcEut respectively and blends having peaks depressed are denoted in the legend by Depr.

59

The other two, Figure 1-22 and Figure 1-23, are typical of those seen in later

chapters, one being a eutectic crystallisation and the other being a

Flory-Huggins style crystallisation. There is some uncertainty in the phase

diagrams having Flory-Huggins crystallisation as to whether the melting

having near-constant melting temperature is a true eutectic or not but the

term eutectic will be used in the text. Figure 1-22 takes the same form as

Figure 1-21 except that Liquid-Liquid phase separation is replaced by

melting and Liquid-Solid phase separation for the crystallisation of the

diluent.

110120130140150160170180190200210220230

0 10 20 30 40 50 60 70 80 90 100Polyamide concentration (%wt)

Tem

pera

ture

(0 C)

TmPA69PureTmPA69DeprTmPThDeprTmEutTcPA69PureTcPA69DeprTcPThDepr

Solid

Liquid

Solid & liquidLiquidLiquid & crystallites

Liquid

SolidSolid & liquid

Solid & liquidLiquid

SolidLiquid & crystallites

Solid & liquid

SolidSolid & liquid

Liquid & crystallites LiquidLiquid

Liquid & crystallites

Figure 1-23 Example of Flory-Huggins style non-equilibrium phase diagrams for heating to the liquid state and cooling raw materials and blends of polyamide-69 (PA69) and carbazole (Car) Melting peaks are noted with Tm and crystallisation by Tc. Blends having peaks depressed are denoted in the legend by Depr.

The eutectic style and Flory-Huggins style crystallisations presented in the

various chapters are consistent with the above two types of phase diagrams

in that samples are from ampoule material and temperatures are peak

temperatures. Red coloured text and graphics refer to heating at 5 0C/min

whilst blue is for cooling at 2 0C/min. Phase regions described in black are

common to heating and cooling. Reference to crystallites is where there have

been a small amount of near pure crystallites of polymer with melting and

crystallisation temperatures very close to the pure polyamide. The phase

regions for them are delineated in the liquid region by fine broken lines.

Heavy long dashed lines delineate the melting/crystallisation of polyamide.

60

Heavy short dashes delineate the melting/crystallisation of diluent. Solid

lines delineate the melting/crystallisation of eutectics. Data points are solid

diamonds for pure polyamide, solid squares for polyamide in blends

depressed in peak temperatures, triangles for the diluent and solid circles for

eutectics.

It can be seen in many of the phase diagrams that, where polyamide and

diluent where there has been Flory-Huggins style crystallisation to be also

defined as melt or crystallise almost simultaneously, there is a slight

depression of the transition temperature in a manner similar to that

described by Berghmans [135]. The crystallisation of diluents giving “spiky”

peaks allows definitive assignment of crystallisation peaks. That is not the

case for melting where peaks are Gaussian in form. The assignment of

melting peaks to polyamide or diluent has often been determined for melting

phase diagrams from the order in which the crystallisation has taken place,

recognising that whichever has the higher crystallisation temperature will

also have the higher melting temperature for the same blend. Account is

also taken with this of the area of the peaks in relation to the amount of

material of each in the blend. This has allowed phase diagrams for melting

in the case Flory-Huggins melting rather than as eutectic melting because

often the first melting peak is taking place at near-constant temperature.

There are three alternatives for phase diagrams. One is to consider the

onset of eutectic melting as the solidus and the end of a TLS peak as the

liquidus but this leaves an unusual gap between them right at the eutectic

point caused by the difference between start and ending of eutectic melting.

Another is to do as Visjager, Tervoort and Smith [136] and combine peak

temperatures of eutectic melting with the end of the TLS peak but this is an

unusual combination and would require a different approach where

Flory-Huggins style melting or crystallisation takes place. It has been

decided for consistency to use peak temperatures throughout the phase

diagrams, giving consistency in presentation of information throughout the

whole thesis. As a caveat, it should be recognised that the experimental

conditions differ strongly from equilibrium and that the use of peak

temperatures is not giving the temperatures at which all material is in the

liquid (or solid) state.

61

1.5.3 Simultaneous Differential Thermal Analysis/Thermogravimetric Analysis

Simultaneous Differential Thermal Analysis/Thermogravimetric Analysis

(SDT) allows combined TGA and DSC to be run on a sample at the same

time. It does not have quite the same TGA or DSC sensitivity as the

individual instruments. It does have major benefits in allowing rapid

assessment of materials from a large selection of materials. Melting,

evaporation and degradation can be assessed from a single fast experimental

run. This allowed the efficient selection of candidate materials for the work

based on having high melting temperatures and the material not evaporating

too quickly in the working range to 300 0C.

A typical analysis curve obtained from this technique is shown below in Figure 1-24.

Figure 1-24 Example of Simultaneous Differential Thermal Analysis

with Thermogravimetric Analysis (SDT) for Flourenone.

The dashed curve of the temperature difference between the sample and an

inert reference shows the melting of fluorenone just above 80 0C followed by

the evaporation of the molten fluorenone. The solid curve does not show any

significant weight loss at the melting temperature but the sample mass has

been reduced to zero by 255 0C due to evaporation in the nitrogen gas

stream. The maximum feasible working temperature for combination with

polyamides in a DSC pan would be just above 150 0C, too low for polyamides

that melt at 209-290 0C.and may have to be taken up to 307 0C to remove

residual lamellar nuclei. Approximately 20 candidate materials were quickly

62

evaluated this way, leaving carbazole and phenothiazine as the only

reasonable ones left.

1.5.4 Fourier Transform Infrared Spectroscopy 1.5.4.1 General

Infrared spectroscopy is a technique that measures the interaction between

a material and infrared (IR) frequency electromagnetic radiation. It is

commonly used in the mid range frequencies of 4000 to 400 cm-1. This is the

region where molecular vibrations and rotations show absorbance bands

that are characteristic of the atoms involved in the bonds. Atomic bonds in

a molecule can generate absorbance bands in the far, mid and near infrared,

regions. Energy is absorbed when the frequency of the irradiating

electromagnetic waves is in resonance with characteristic modes of

molecular movements such as bond stretching, vibrations and rotations.

Changes in the states of groups of atoms in a polymer molecule can cause

shifts in the absorbance bands. These can contribute to our understanding

of what is happening to the molecules or parts of molecules and it is this

area that is important when looking for the influence of changed hydrogen

bond environments [42, 112, 113]. Each specific chemical environment will

have characteristic frequencies where infrared radiation is absorbed. In

principle, this allows the determination of molecular structures from the

infrared “fingerprint”, however full interpretation can be difficult because of

the myriad of different possibilities with reasonably sized organic molecules.

It is quite easy, for example, to discriminate at four or five places across the

mid IR spectrum between the various polyamides used in the trials. Various

authors have discriminated between amorphous and crystalline states of

polyamides [82] and between various crystallographic forms of polyamides

[82, 137, 138] including Brill transitions [91] using a variety of FTIR

techniques. Polyamide interactions with liquid crystal oligomers have also

been detected [139], as have those with fibre reinforcement [140].

The majority of work on Fourier transform infrared spectroscopy (FTIR) is in

the mid range of the IR spectrum but the Near IR (NIR) in the range 11,000

to 4000 cm-1 is very sensitive to subtle differences in hydrogen bonding.

We will see later that the original premise of hydrogen bond destruction with

polyamides by the potential hydrogen bond acceptors, carbazole and

phenothiazine, was cast into doubt by the mid range IR work. This prompted

63

validation by NIR investigations in general bands identified as being related

to hydrogen bond interaction with polyamides [141].

Earlier instruments were dispersive but these have largely been supplanted

by inexpensive, rugged Fourier transform infrared spectroscopy instruments

that allow relatively quick measurements to be made with extremely high

signal-to-noise ratios. A Nicolet 750 FTIR instrument was used for the

work.

There are a variety of FTIR techniques available to use:

a) Attenuated Total Reflection(ATR)

b) Transmission of solutions

c) Transmission of cast thin films

d) Diffuse Reflectance Infrared Fourier Transform (DRIFT) spectroscopy

where the infrared sample beam is deflected downwards onto the surface

of a sample with an elliptic mirror. Any diffuse IR reflection from the

surface is collected with another elliptical mirror and focussed back into

a beam incident on the detector. One advantage of this technique is that

samples directly as formed may be examined without disrupting the

morphology or chemical interactions between molecules.

e) Photoacoustic Spectroscopy (PAS) studies utilise the generation of

thermal waves in the sample upon infrared absorption. This leads to

acoustic waves being propagated within the sample and into the

surrounding gas. A sensitive microphone picks up the acoustic signal

and amplifies it to give spectra as the IR frequency is swept across the

mid infrared spectrum. The original principle dates back to the 1800s.

It has been applied in FTIR for the last dozen years or so.

The first three techniques were not utilised because they involved

modification of the bulk samples in ways that would alter their morphology

at the detecting surface or in the bulk.

Photoacoustic (PAS) detection was used for the Mid IR range experiments

because it is more suitable than DRIFT when looking at small differences in

frequency. The photoacoustic approach does have a disadvantage in that

the heights and areas of peaks do not necessarily represent the relative

intensities of the absorption of IR. There can be an attenuation of strong

64

signals. This will be discussed below. In this particular case, the

advantages of using material in its native morphology and having very

accurate peak frequencies outweigh the drawbacks due to non-linearity.

DRIFT was used for the NIR experiments because the instrument signal was

far superior to the photoacoustic signal with that part of the IR spectrum.

1.5.4.2 Mid Range IR and hydrogen bond Interactions

Polyamides have an N-H stretch with a large peak near 3300 cm-1. The

normal situation for polyamides is to be strongly hydrogen bonded from the

carbonyl oxygen through the amide hydrogen to the nitrogen of another

amide group on the same polymer molecule or another molecule. The large

peak near 3300 cm-1 represents the bound state because the vast majority of

potential hydrogen bonds are consummated at room temperature [48

p. 270].

The state of the N-H bond in carbazole material is normally unbound. There

is a major N-H peak at 3441 cm-1. There should be a shift in the IR peaks

for the N-H from polyamide-4,6 and the N-H peak from the carbazole if the

carbazole molecules replace polyamide N-H in the hydrogen bond structure.

The carbazole N-H will then become bound and the polyamide N-H will

become unbound. There should have been shifts in both towards each other

of about 10 cm-1 if there were any substantial complexing of the two

materials with hydrogen bonding.

Guerra et al. [142] found shifts of 58 cm-1 in N-H stretching band maxima as

they altered the percentages in their hydrogen bond interacting blends.

The N-H stretch for polyamide-6 film increases by 18 cm-1 in being heated

from 50 to 227.5 0C in work by Xu et al. [143] due to the reduction in bound

hydrogen bonds and a move to less restricted N-H bonds. The same paper

shows a shift to the right in the melt of polyamide-6/LiBr compared with

pure polyamide-6 because the amide-amide hydrogen bonds are supplanted

by the intense ionic bonds with the salt.

Gao and Scheinbeim studied interactions between Nylon-11 and

poly(vinylidene fluoride) (PVF2) [144]. They found a shift in the N-H stretch

by up to 8 cm-1 as the level of PVF2 was increased. This shift was to lower

wavenumbers because the F…N-H hydrogen bond was stronger than the

C=O…N-H bond. That is obviously in the opposite direction to that expected

65

if the N-H of carbazole or phenothiazine were to supplant the amide N-H

bond to O=C on another amide group thus freeing up an amide N-H. It is

therefore a useful benchmark for the type of change expected.

Skrovanek et al. also looked at semicrystalline Nylon-11 considering the

effects of temperature increases leading to the melt [61]. They found a shift

in the peak of the main N-H stretch of 32 cm-1 to higher wavenumbers in

that process as the temperature was raised and the hydrogen bonds

weakened. In their case the normalised area of the peaks reduced in

sympathy with the temperature increase.

Wang, Ma and Wu [145] solution blended polyamide-6 or polyamide-6,6 into

“Novolac”, a phenolic resin. The aim was to reduce the brittleness of the

Novolac by using intermolecular hydrogen bonding of the materials in the

blends. They did not specify explicitly in the paper the extent of the FTIR

frequency shifts they found but their figures 5 and 6 plotting the spectra for

various blends make it clear that substantial shifts have, in fact, taken

place. The O-H of the Novolac has changed by something in the order of

70 cm-1.

The focus of this Mid-Range IR work will be on the N-H stretch as that is the

major area where the disruption of C=O….H-N(amide) with “free” N-H(diluent)

to produce C=O….H-N (diluent,-“bound”) and “free” N-H(amide) would be

expected to have an effect. Peak frequency was used as the determinant of

N-H changes rather than the more risky deconvolution of non-linear PAS

signals of composite spectra from different materials (vide infra). The

results, above, from other authors’ work gave the confidence that this would

be suitable to discriminate changes in hydrogen bonding activity.

1.5.4.3 Mid Range IR Frequencies of Interest

The relevant absorption frequencies for FTIR investigations described in

Chapters 3 to 10 for the different combinations of polyamide-4,6

polyamide-6, polyamide-6,9 or polyamide-6,12 with carbazole and with

phenothiazine are brought together in Appendix C: FTIR Assignments to

avoid undue repetition.

The Photoacoustic (PAS) spectrum of polyamide-4,6 is seen in Figure 1-25

below. The other polyamides in the study have peaks that are close but not

66

quite identical to the above. The slight differences across several absorbing

bands can be used to positively identify polyamide types.

Polyamide-4,6

5

10

15

20

25

30

35

40

45

50

55

60

65 Ph

otoa

cous

tic

500 1000 1500 2000 2500 3000 3500 Wavenumbers (cm-1)

Figure 1-25 Mid Range IR spectrum of polyamide-4,6 from an ampoule

The bands for carbazole are shown in the PAS spectrum of Figure 1-26 and

for phenothiazine are shown in Figure 1-27.

Carbazole

10

20

30

40

50

Phot

oaco

ustic

500 1000 1500 2000 2500 3000 3500Wavenumbers (cm-1)

Figure 1-26 Carbazole photoacoustic FTIR peaks in the Mid Range IR

67

Phenothiazine

40

80

120

160

200

Phot

oaco

ustic

500 10001500 2000 2500 3000 3500Wavenumbers (cm-1)

Figure 1-27 Phenothiazine photoacoustic FTIR peaks in the Mid Range IR.

1.5.4.4 Mid Infrared Data Analysis for Blends

The original PAS spectra of polyamide-4,6 (Ampoule 64) and carbazole

(Ampoule 63) are overlaid in Figure 1-28 with the spectrum of the ampoule

material 66PA46Car (Ampoule 31) to demonstrate an FTIR analysis problem.

Carbazole Polyamide-4,6 66PA46Car

5

10 15

20

25

30

35

40

45 50

55 60

65

Phot

oaco

ustic

500

1000

1500

2000

2500

3000

3500

4000

Wavenumbers (cm-1) Figure 1-28 PAS spectra in Mid IR for carbazole, polyamide-4,6 and a blend.

It can be seen from Figure 1-28 that each of the spectra for the raw

materials has a large number of sharp peaks. The spectrum for ampoule

68

material from a blend takes on approximately the combined peaks of the two

raw materials. The peak from one constituent material of a blend may lie on

a sharply rising or falling portion of the other material’s spectrum. The

combined effect can result in a shift in peak frequency even if there are no

changes in hydrogen bond interactions or morphology due to blending. The

Photoacoustic technique is usually non-linear for strong peaks. Spectral

additivity cannot normally be expected. There was a conundrum, however,

because there was a large problem here with the interpretation of spectra.

The following mathematical modelling was employed in an attempt to see if

the infrared spectra indicated interaction between the polyamide and the

small molecules. The spectra of the two constituents were mathematically

added in a proportion that mimicked the salient features of the spectrum of

the blend material. It was done in order to look for regions where the blend

spectrum was different from that expected for no interactions involved.

Differences between the model and experimental results could potentially be

indicative of frequency shifts. It was reasonably strong evidence for no

hydrogen bond interaction or crystallographic/morphology changes to have

taken place if the “model” and experimental peaks matched up precisely.

Any artefact caused by the simple model would have to be exactly the same

magnitude but of the opposite direction to actual chemical shifts, an unlikely

scenario. The importance of relative height changes of double and treble

peaks was considered low.

Regions of each constituent spectrum were chosen where there was a

significant peak in one material but not in the other. The spectra of the two

materials were mathematically added and scaled to match the spectrum of

the ampoule material at both points. Sometimes more points were selected

to assist in the match. Generally the peaks that were chosen were ones

where the signal for one material was reasonably high and the other material

had a low signal at that point. The highest peaks were not chosen as primary

ones as they were likely to be non-linear due to sensor saturation. The

match to the actual identified peaks would thus be due to having the correct

proportions of each material in the model. Examples of the spectral regions

chosen are given in Figure 1-29. Secondary peaks have the spectrum of the

other material is rising or falling strongly in that region or the peak is a very

high one which is likely to be truncated by signal saturation.

69

Carbazole peak and low Polyamide-4,6 absorbanceCarbazole confirmatory peak

Confirmatory peak for Polyamide-4,6

Polyamide-4,6 peak andlow carbazole absorbance

Carbazole Polyamide-4,6

2

4

6

8

10

12

14

16

18

20

22

24

26

28

30

Phot

oaco

ustic

5001000150020002500300035004000Wavenumbers (cm-1)

Figure 1-29 Peaks used for modelling polyamide-4,6/carbazole blend spectra in Mid IR.

An example of a model compared with a measured spectrum for a blend is

given in Figure 1-30.

*Addition* model 23PA46Car

4

6

8

10

12

14

16

18

20

22

24

26

28

Phot

oaco

ustic

5001000150020002500300035004000Wavenumbers (cm-1)

Figure 1-30 Mid IR PAS spectrum of 23PA46Car from Ampoule 57 with the model constructed from the spectra from polyamide-4,6 and carbazole.

This is expanded for the carbazole N-H stretching peak and the peak heights

equalised in Figure 1-31.

70

*Addition* model 23PA46Car

12

14

16

18

20

22

24

26

Phot

oaco

ustic

33803390340034103420343034403450Wavenumbers (cm-1)

Figure 1-31 Carbazole N-H stretch for 23PA46Car and model expanded from the comparison in Figure 1-30 and equalised in height.

PAS does have a disadvantage that strong peaks will cause signal saturation

resulting in some truncation of peak heights. That applies particularly to

the major peaks for spectra of the constituent materials. The

mathematically combined spectra will therefore also give partly compressed

tops of major peaks. These peaks are then expanded vertically in the graphs

to give peaks that can be easily compared for peak frequency with the

measured ones for the polyamide/diluent blend material from the ampoules.

The rounding of the peak tips, as seen in Figure 1-31, is an artefact of

expanding the models based on non-linear spectra to give the same peak

heights for comparison. The reason the models are built primarily on

significant peaks that are not the largest peaks for each material is to

construct the model minimising the effects of signal saturation. The spectra

for polyamide/diluent blend samples are less compressed at peak tops

compared with the constituent spectra that the models were based on. Each

highly absorbing peak is less saturated in the measured blend spectrum.

Peak signals for blend materials are reduced in intensity by dilution because

of the presence of the other material in the sample. The peak signals are

then more linear because there is less detector saturation. The positions of

peaks are the critical issues rather than the rounding.

71

It can be seen above that analysis without mathematical modelling would be

extremely difficult but simple mathematical addition of spectra can result in

artefacts due to non-linearity in the major spectral peaks for the constituent

materials. A comparison between a measured blend spectral peak and a

model from spectra of the constituent materials with no change in the peak

frequency should be strong evidence for the materials not interacting. Slight

height differences between sharp measured peaks and a model based on the

spectra of the constituents may or may not be indicative of an interaction,

given the non-linearity of the detector system.

1.5.4.5 Near Infrared FTIR (NIR)

The NIR region is noted in the literature [141] for being sensitive to the

hydrogen bond status. The broad FTIR peaks for hydrogen bonding can be

examined in the Near Infrared (NIR) at moderate sensitivity and resolution.

The area of interest is in the 7500 to 4000 cm-1 region. The Nicolet 750

instrument can be set up for the Near Infrared region with wavenumbers

between 11,000 and 3,000 cm-1. This requires some changes to the physical

configuration of the instrument regarding the beam splitter, light source and

detector. The DRIFT technique was the most appropriate.

Wu and Siesler [141] studied polyamide-11 in this band. The frequencies in

the Near Infrared they found were 6912, 6600, 6390, 6290, 6180, 4940,

4846, 4580 and 4560 cm-1. Values found with the polyamides here were

close to those values. A typical polyamide spectrum is shown in Figure 1-32.

Abso

rban

ce

4500 7500 Wavenumbers (cm-1) Figure 1-32 Typical polyamide DRIFT spectrum in the NIR region.

No NIR peak values were located in the literature for carbazole or

phenothiazine.

72

The spectra for these two are shown below in Figure 1-33

carbazole phenothiazine

Abso

rban

ce

4500 7500 Wavenumbers (cm-1) Figure 1-33 Carbazole and phenothiazine Near Infrared spectra as

measured with DRIFT.

The peaks for the polyamides compared to those of the diluents are quite

separate and are not so steep as in parts of the Mid Range IR so it was not

necessary to resort to the mathematical additions of spectra. There was a

facility in the Omnic software used to drag an individual spectrum from a

group up or down to match the height on screen of another spectrum peak.

That facility was used to move each peak from the spectrum for a blend to

match the equivalent peaks for those of the spectra for the constituents.

1.5.5 Small Angle X-ray Scattering Small Angle X-ray Scattering (SAXS) can be used to gain information about

semicrystalline polymers at a dimensional range around the size of lamellae

and relates to their stacking within a solid. Periodically stacked lamellae

reflect X-rays more strongly than amorphous regions because of the higher

density (and therefore electron density). This occurs with Bragg reflections

at scattering angles of 2Θ according to the well known relationship of sin

Θ = mλ/2d where m is an integer giving the order of scattering, λ is the

wavelength of the X-rays (in this case) and d is the periodicity distance of the

scatterers. For example, an X-ray wavelength of 0.154 nm for first order

scattering giving a peak at 2Θ of 0.80 equates to a scattering periodicity near

12 nm.

An opportunity arose early in the making of ampoules to have some samples

measured with SAXS by another University of South Australia PhD student

who was working briefly at Connecticut University. He was able to carry out

a very restricted number of trials because, at that stage, only a handful of

73

ampoules had been produced. The conclusions from these are limited but

the results have been included for the benefit of other researchers.

1.5.6 Solid state Nuclear Magnetic Resonance Spectroscopy FTIR, discussed above, looks at how the frequencies of interatomic bonds are

influenced in their various modes of vibration by the atoms at the ends of

the bonds and by the near-neighbour environment. Nuclear Magnetic

Resonance (NMR) spectroscopy is interested in the nuclei of atoms and how

they are influenced by near-neighbours. The nuclei of 1H, 13C 15N and some

isotopes of other elements act as if they are spinning. It is possible, by

placing them within strong magnetic fields and irradiating them with radio

frequency (rf) electromagnetic radiation, to have them absorb energy. They

appear to precess in the way the top of a spinning top gyrates in slower

circles as the top spins. The combinations of magnetic field and rf frequency

where the absorption occurs can be used to determine the environments of

the nuclei. For example, the five hydrogen atoms attached directly to a

benzene ring having an attached O-H group will have absorptions at three

slightly different frequencies, one for the H atom directly opposite the O-H

bond, there will be a peak twice as large and at a slightly different frequency

for the two atoms next to it and another peak the same size as the previous

one for the two hydrogen atoms attached to carbon atoms either side of the

O-H group. There will also be a peak of the same size as the original peak,

but at a noticeably different frequency, for the hydrogen atom attached to

the oxygen atom of the O-H group. Similarly differing 13C environments will

produce different 13C peaks depending on the environments of the nuclei.

The measurements can either be carried out by keeping the rf irradiation

constant and modulating the strong magnetic field or by keeping the

magnetic field constant and varying the rf field. Usually the latter approach

is taken nowadays with the advent of strong superconducting magnets and

the ability to give a sharp pulse of rf radiation which populates all the atoms

at the same time. The response as a function of time is deconvoluted to give

the final output.

These measurements can be carried out in solution or of solid materials.

Solidified small organic molecules have sharp peaks because the

environments around the nuclei are quite regular. Even highly crystalline

polymers always have a considerable amount of disordered amorphous

74

polymer material outside the spherulitic regions and in the interlamellar

spaces. This leads to the absorption peaks being smeared out. It is often

possible, however, to infer things from solid state NMR spectroscopy where

the material is being left in its original morphological state.

It was for this reason, that the opportunity to have solid state NMR

measurements made on some blend material from ampoules was taken up

in order to better understand the materials we were working with. A brief

window in time occurred after the production of Ampoule 1 for NMR

measurements to be carried out by Dr. Andrew Whittaker at Queensland

University on red and white sections from the ampoule plus polyamide-4,6

and carbazole powders. The white sample proved too hard to make into a

fine powder at the time so only the red blend could be measured along with

the constituent powders. The results proved ambiguous. They have been

included in Chapter 3 to make them available for other researchers.

1.6 Structure of the Thesis The thesis covers much work that is of the same structure from chapter to

chapter covering different material combinations. A brief description of the

various chapters follows.

Chapter 2: Experimental

The second chapter covers all experimental details and scant reference is

made in other chapters to these details.

Chapters 3 to 10: The polyamides combined with the diluents

These chapters cover the experimental work on various combinations of

polyamides with either carbazole or phenothiazine. A chapter is devoted to

each combination. The polyamides polyamide-4,6, polyamide-6,

polyamide-6,9 and polyamide-6,12 in that order are combined firstly with

carbazole. The last four chapters of this block cover the above four

polyamides combined with phenothiazine.

There is much that is similar from chapter to chapter within this group of

chapters but there is also much information that is different from one

material combination to another. As a result of the consistent approach

taken in the experimental work and in analysing the data, the chapters

could appear repetitive. This will not be helped where the outcomes from

one material combination to the next happen to be similar.

75

The chapters give only low-level conclusions on the results seen. This is the

most appropriate because often the outcomes of experimental work on one

or more combinations of the various polyamides with either of carbazole or

phenothiazine are best evaluated together in the final conclusions chapter.

Chapter 11 General Conclusions

This chapter draws the previous eight chapters together in overall

conclusions. The differences between the polyamides in the way they interact

in high temperature solution and in crystallising to solid blends are covered

in the context of the molecular structure of the individual polyamides. The

common aspects are more oriented to how polyamides generally interact with

these specific small molecules.

This work has covered a reasonable tract with non-isothermal DSC and FTIR

and could be considered as a pilot study. There are a number of aspects

that could be pursued to further the scientific understanding and to pursue

applications of this research. A list of questions that the work raises is

provided in the hope that the opportunity will arise for them to be

investigated.

Appendices

Appendix A: Further details from DSC thermograms

Appendix B: Lissajous Figures for Understanding Temperature Modulated

Differential Scanning Calorimetry of Nylons

Appendix C: Mid Range Fourier Transform Infrared Spectroscopy

Assignments

Appendix D on CD: Fourier Transform Infrared spectra (PDF format) of

blends with mathematical models or with spectra of

constituent materials. The CD contains the whole thesis

Bibliography

Bibliography of all references in the thesis.

1.7 Summary This introductory chapter to the thesis explained how aliphatic polyamides,

commonly called nylons, are an important class of engineering polymers,

that it is important to understand their properties more fully to utilise them

76

to best advantage and how this work contributes to the virtually untapped

knowledge of their characteristics in high temperature solutions with low

molecular mass diluents. It went on to take the reader through from the

background information on mixtures of materials, hydrogen bonding, the

structure, melting and crystallisation of semicrystalline polymers, and led to

the specific case of aliphatic polyamides. This continued into a survey of

some of the recent literature in areas adjacent to the specific area of interest,

leading to a description of the research problem. A description was then

given of techniques suited to investigating the research problem and the

reasons. The details of expected outcomes of certain techniques were

provided in some cases including the “TLS peak” often found when heating

blends in the DSC and of mathematical modelling blend spectra in the Mid

IR range using the spectra of the constituent materials.

It is now time to look at Chapter 2 with its description of the experimental

conditions used for the techniques.